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Contrasting points from Plansee
technical trends
Contrasting points from Plansee Ken Brookes reports on two very different subjects of considerable practical importance covered at the Plansee hardmetals conference: the sinter bonding of steel and hardmetal and the suppression of WC/Co recrystallisation and grain growth without specific inhibitors...
T
he novel technique for the fabrication of steel/hardmetal composites, presented by Arno Huber of Robert Bosch GmbH’s Corporate Sector Research and Advance Engineering, GerlingenSchillerhöhe, Germany, attracted attention at Plansee. Sinter bonding combines densification of the steel part and its joining to fully sintered hardmetal in a single step. They cannot be sintered simultaneously because the steel would melt at the carbide’s sintering temperature, but substantial interdiffusion at the sintering temperature of the steel achieves a strong atomic bond. Although the basic principle is fairly obvious, there are major problems in commercialising the concept. Most important of these is the differential shrinkage of the two components, causing warpage, high internal stresses or even fracture during the cooling phase from sintering temperature. In brazing, carried out at substantially lower temperatures, the similar effects can be combated by absorption of stress by the braze metal or deformable gauze. In the research discussed in this paper, the carbide component fits loosely (though to a carefully calculated degree) in the steel cavity at sintering temperature, the pressure needed for diffusion bonding being applied by steel shrinkage during cooling. Shrinkage stresses thus augment the sinter bonding. High-carbon steels like 100Cr6 work well in this application. By contrast,
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MPR February 2010
low-carbon steels such as 17-4PH or maraging steels absorb carbon from the carbide constituent of the hardmetal and then form a brittle interlayer consisting of Ȗphase, adversely affecting the mechanical properties of the composite. The microstructural and mechanical characteristics of the bonding zone in several combinations of steels and hardmetals were investigated for this research. Hardmetals, also known as cemented or sintered carbides, exhibit excellent wear resistance, hardness up to around 2000 HV, high transverse rupture strength, high Young`s modulus and compressive strength up to 6000 MPa. However, hardmetals are comparatively brittle, thus prone to breakage caused by shear and tensile stresses, and their material and production costs are high compared to steel. Hence it is common practice to join hardmetal parts to steel supports, typically by brazing, welding, diffusion bonding, overlaying or the use of mechanical methods or adhesives. Some important aspects of different joining methods are listed in Table 1. To study the applicability of sinter bonding to different steel/hardmetal combinations, two different hardmetal grades and four steel grades, listed in Table 2, were selected for the experiments. The hardmetal WC/10Co is a very common grade used in wear applications while the Ni-containing grade has high compression strength and better corrosion resistance,
making it suitable for high-pressure applications. The basic criterion for the steel selection was high yield strength, as the support for hardmetal should be strong and rigid. The steel 100Cr6 exhibits high strength, hardness and wear resistance and is often used for bearings. Where corrosion resistance plays a role, X65Cr13 may substitute for 100Cr6. The high-alloy maraging steel 18Ni350 combines tensile strength of about 2450 MPa with high ductility. Precipitates like Ni3Ti and Ni3Al are responsible. 17-4PH is also a highly alloyed martensitic precipitation-hardening stainless steel, commonly produced by powder injection moulding (PIM) because of its easy sintering and heat treatment. For 100Cr6, two powders with different grain size were used. The finer powder with d50 5.7 μm was used in combination with N09 and the other with OM10. The most important characteristics of the selected steels and hardmetals are listed in Table 2. The sinter bonding experiments employed cylindrical hardmetal inlays and steel green parts with a corresponding cavity. To provide suitable hardmetal inlays, OM10 powder feedstock was prepared for PIM. For this purpose, 59 vol% of hardmetal powder was premixed with an organic binder containing about twothirds polyamide and one-third paraffin wax. This feedstock premix was homogenised and granulated in a twin-screw
0026-0657/10 ©2010 Elsevier Ltd. All rights reserved.
Table 1. Overview of common methods for joining steel and hardmetal. Joining method
Advantages
Disadvantages
Max. shear strength
Brazing
- high reliability - cost and time efficiency
- high residual stresses - insufficient wetting may occur - hazard of toxic fumes caused by solder and fluxing agents
320 MPa
Welding (e.g. electron beam welding, friction welding)
- thermal stresses are minimized compared to brazing - applicable for large parts
- brittle γ-phases, pores and cracks may form in the joint zone
Comparable to brazed joints
Adhesives
- no thermal stresses are induced - bonding zone is thermally and chemically not affected - health hazards and capital costs are low
30 MPa - low strength - only adequate for low operating temperatures
Mechanical attachments (clamping, wedging, screw mounting)
- easy substitution of worn parts - no thermal stresses are generated - retention of optimum heat treatment properties
- no substance-to-substance bond - stress concentrations are induced - geometric limitations
Shrink fit mounting
- homogeneous stress distribution - no substance-to-substance bond - temperature-dependent bond - compressive stresses are strength induced within the hardmetal (pre-load)
Diffusion bonding
- thermal stresses are partly removed by a ductile interlayer - high strength - joining areas may be very large
compounder. Cylindrical bars were injection moulded with a diameter of 6.0mm and cut into 25mm lengths. The green samples were machined to a pencil-like shape with cone angle of 120° and diameter 5.6-5.7mm. Afterwards the samples were cut to a length of about 6mm and the paraffin wax binder mainly removed in acetone (65°C, 1000 min). Thermal debinding was completed and sintering
- mating surfaces must be perfectly planar and clean - time and cost consuming
carried out in a single one-hour hydrogenatmosphere stage up to 1400°C in a Gero graphite furnace. To increase carbon activity in the atmosphere but minimise interdiffusion of carbon, the samples were placed on alumina powder in a closed graphite boat. As Ni-based hardmetals had to be processed in vacuo, sintered inlays of N09 with similar geometry were purchased from Tribo GmbH. In all sintered samples
Depending on geometry and joining method
Depending on oversize
380 MPa
both free graphite and eta phase were absent. The steel feedstocks were prepared in a similar manner, but a powder loading between 59 and 68 vol% was chosen, depending on the powder characteristics. Plate-like specimens were injectionmoulded with a small borehole placed in each centre. A wider, concentric blind hole with a cone angle of 120° was drilled, to serve as a seat for the hardmetal inlay. The
Table 2. Specifications of selected hardmetals and steels. Designation
Composition1 wt%
Grain size2 d50 µm
Strength3 MPa
2.5
not specified
Multi Metals
1
≥ 3,600
Tribo Hartstoff GmbH
5.7 / 8.2
1,600 – 2,400
0.4
Sandvik Osprey Ltd
Elongation %
Supplier
Hardmetals: OM10
WC/10Co
N09
WC/8Ni
Steels: 100Cr6 (1.3505)
Fe-1C-1.6Cr-0.35Mn-0.23Si
X65Cr13 (1.4037) Fe-0.63C-13.8Cr-0.76Mn-0.12N
8.1
570 (cast)
not specified
Sandvik Osprey Ltd
18Ni350
Fe-17.5Ni-12Co-3.3Mo-1.2Ti0.1Mn-0.1Cr-0.08Al
4.2
2,400
6
Sandvik Osprey Ltd
17-4PH (AISI 630)
Fe-17Cr-4.1Cu-4Ni-0.48Mn0.36Si-0.23Nb-0.17Mo-0.11N
8.6
515 – 1,170
10 – 18
Sandvik Osprey Ltd
1 Manufacturer’s data 2 Hardmetals: grain size after sintering; steels: grain size of the powder 3 Hardmetals: transverse-rupture-strength (manufacturer’s data); steels: yield strength σ0.2 of wrought material (literature data)
metal-powder.net
February 2010 MPR
15
steel parts were also debound in acetone at 65°C for 1000 minutes. For the sinter bonding process, the hardmetal inlays were inserted into the cavity of the steel green body and the samples sintered in a hydrogen atmosphere. Figure 1 indicates that the gap between steel and hardmetal was closed during sintering, caused by the shrinkage of the steel part, and a strong atomic bond achieved due to interdiffusion. The author explained that the diameter of the cavity had to be chosen so that an apparent oversize of the hardmetal inlay was created during sintering. In this way the surfaces came into close contact, supporting the formation of the bond. The oversize would not occur at the sintering temperature, since diffusional flow or creep would deform the steel part as soon as a contact was established. The apparent oversize could be determined accurately by measuring the total shrinkage of reference samples of the respective steel. The hardmetal inlay had to be able to slide upwards during the sintering process, indicated by the arrows in Figure 1. If sliding were supressed, by too high an inlay weight, for example, bonding at the side surfaces would be poor and the steel part might deform. Thus bigger and heavier parts should be sinter-bonded using a “mounting rack” of some kind. Some of the sintered composites were heat-treated in order to study the influence of interdiffusion between the two components on the microstructure and mechanical properties. The samples were studied with both optical and scanning electron microscopy. Element distribution was analysed by EDS (energy dispersive X-ray spectroscopy). Vickers hardness was measured with a 10kg load and microhardness across the boundary surface, to study the influence of interdiffusion on the mechanical properties of the heat-treated samples. As the hardness values differed heavily the microhardness load was varied between 25g and 500g. As expected, the quality of the steel/ hardmetal bond depended strongly on the apparent oversize of the hardmetal inlay after sintering. Higher oversize resulted in earlier and closer contact between the mating surfaces, which in turn promoted diffusional flow. However, if the oversize were too high, noticeable deformation of the steel part might occur. Optimal oversize depended on the fit between the two parts, their surface finish and the materials. Oversize values between 1% and 5%
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Figure 1. Schematic illustration of the sinter bonding process. The sintered hardmetal inlay is inserted into the cavity of the steel support green body. The gap between both parts is closed upon sintering and a strong bond develops due to interdiffusion.
seemed reasonable. Diffusion mechanisms played a crucial role in bonding and oxide layers could significantly impede interdiffusion. Fortunately, oxides could generally be removed at an early stage of sintering by applying a reducing hydrogen atmosphere or vacuum, avoiding costly cleaning of the mating surfaces. Due to interdiffusion between the steel and hardmetal binder phase, a solid solution occurs first at the interface and then intermetallic compounds or precipitates depending on the compositions of both materials. In this regard it was advantageous that Fe, Co and Ni exhibited face-centred cubic (fcc) structures and were completely intersoluble at sintering temperature. The sintered 100Cr6-WC/10Co composite samples showed excellent bonding between steel and hardmetal inlay, illustrated in Figure 2A. A narrow porous layer formed near the interface, indicating strong interdiffusion between the parts. Adjacent to the interface, a bright, 300 – 400μm thick, zone was visible in the hardmetal component. During cooling, high residual stresses were generated in the composite due to the significant difference between the coefficients of thermal expansion (CTE) of steel and hardmetal. The greater contraction of steel led to high compressive stresses in the hardmetal inlay which was beneficial for the durability of the assembly. First, the
compressive stresses generated a clamping force which counteracted potential failure of the bonding zone. Second, tensile stresses in the hardmetal inlay, which could be crucial and might have arisen from the high internal pressure, were partly compensated. However, if the composite was cut axially the stress state was altered and cracks formed in the hardmetal inlay, as shown in Figure 2B. The cracks did not proceed along the boundary but within the hardmetal part, demonstrating high strength in the bond. 100Cr6 had a fully pearlitic microstructure and a bright and diffuse interlayer between both parts, shown in Figure 2C. EDS revealed that this layer predominantly contained cobalt and iron (Figure 3A). The penetration depth of Co and Fe into the steel and hardmetal sides respectively was about 70 – 80μm. In this line scan no diffusion of W could be observed, but the comparison of spectra of different EDS maps showed that some W was present in the Co-rich interlayer. Figure 3B shows a simplified phase diagram of the Fe/Co/Ni system. Co and Ni both stabilise Ȗ-Fe (gamma iron) with fcc structure and austenite is retained at room temperature if these elements are present in sufficiently high amounts. Martensite is formed only in a narrow range near the Ȗ-phase boundary upon cooling. W, Cr and C, also present within the interlayer,
Figure 2. Optical micrographs of the microstructure at the boundary between 100Cr6 and WC/10Co after sintering. The samples were sintered at 1270°C for two hours under hydrogen atmosphere and 100Cr6 showed relative density of 95.8%. A: horizontal section. B: axial section. C: microstructure etched with V2A-solution (room temperature, 180s).
metal-powder.net
decrease the martensite transformation temperature Ms and also the martensite content. This is in agreement with microscopic examinations as no martensite was observed. However, no distinct phase boundary was visible at the interface as would be expected at a pearlite/γ-phase boundary. Instead a smooth transition of pearlitic to a more homogeneous structure was observed, showing small precipitates containing tungsten. This implied that the interlayer had a body-centred cubic (bcc) structure.
Cobalt enrichment at the interface was not expected, as concentration or activity gradients should have been eliminated by diffusion. To study this effect, hardmetal inlays were heat-treated under the same sintering conditions as the composites, avoiding the influence of interdiffusion between steel and hardmetal. After sintering, a 10 – 30μm thick surface layer had formed, most pronounced at the edges, indicating the influence of the sintering atmosphere (Figure 4A).
Figure 3. A: EDS analysis of the element distribution across the interface between 100Cr6 and WC/10Co (1270°C, 2 h). The interface line indicates the borderline of WC grains. In order to reduce scattering each data point was derived from an EDS analysis map by averaging a 35 μm broad zone perpendicular to the line axis. B: Fe-Co-Ni system showing composition range and temperature of martensite transformation.
Hardmetals tend to form so called Ȗphases (M6C or M12C, generally double carbides of the hard constituent and the binder alloy) if the carbon content falls slightly below the stoichiometric value, illustrated in Figure 4B. These phases, brittle and therefore undesirable, can be seen after etching with Murakami solution (equal parts of NaOH and K4[Fe(CN)6]). In Figure 4A the surface layer contains both Ȗphases and a second phase identified as Co binder. It appeared that the carbon-poor atmosphere was responsible for the formation of a Co-rich surface layer, not observed under common (carbon-rich) sintering conditions for hardmetals. Diffusion of cobalt into regions of lower carbon content in hardmetals had been verified by several previous studies. Hence it was assumed that the combined interaction with the atmosphere and the steel was responsible for Co enrichment at the interface. Sintered composite samples were etched to show Ȗphases at the interface, but at most only a minor fraction of dispersed Ȗphase was observed, both in the Co-rich interlayer and in the hardmetal close to the interface. This was surprising, as carbon diffuses interstitially and was expected to be orders of magnitude faster than substitutionally diffusing elements, leading to high diffusional exchange as soon as contact between steel and hardmetal was established.
Higher carbon content
Figure 4. A: Hardmetal inlay showing a surface layer which consists of Co binder phase (bright) and γ-phase (dark). The sintered sample was tempered at 1270°C for two hours under a hydrogen atmosphere, polished and etched with Murakami solution at room temperature for 10 seconds. B: Isothermal section of the W-C-Co phase diagram at 1350°C.
Figure 5. Microstructure at the bonding zone of different steel/hardmetal-composites showing a brittle interlayer. A: Optical micrograph of X65Cr13-WC/10Co (1270°C, two hours). B: SEM micrograph of 18Ni350-WC/10Co (1350°C, one hour). C: SEM-micrograph of 18Ni350 near the interlayer with different precipitates at the grain boundaries, indicated by the arrows (1350°C, one hour).
metal-powder.net
In contrast to 100Cr6, X65Cr13 and the nearly carbon-free steels 18Ni350 and 17-4PH showed a brittle interlayer between steel and hardmetal, seen in Figure 5. This layer mainly contained Co, Fe, W and C as the brittle Ȗphase. Because X65Cr13 contained substantially higher carbon content than the other two steels, interdiffusion of carbon was less pronounced, resulting in a thinner interlayer. Small cracks were visible in the interlayer, due to high residual stresses and precipitates appeared on the steel side near the interface because of interdiffusion. 18Ni350 in particular showed precipitates at the grain boundaries (Figure 5C), which acted as nucleation sites or as preferred diffusion paths for carbon. The bright precipitates contained mainly Mo, W and C and the darker ones Ti and C. It was assumed that these precipitates were carbides. In addition to WC/10Co, a hardmetal with a nickel binder was studied in combination with 100Cr6. A layer consisting of
February 2010 MPR
17
Ni binder and Ȗphase was observed on the free surface of the hardmetal inlays after sintering. A relatively broad interdiffusion zone formed between steel and hardmetal, as shown in Figure 6A, subdivided into two separate layers 1 and 2. As with the Co-based hardmetal, a binder-rich interlayer was formed, due to the carbon gradient from the carbon-deficient interface. At the sintering temperature of 1270°C, iron and nickel form a solid solution with fcc structure over the whole concentration range. However, at room temperature several phases are thermodynamically stable: γ-Ni and the intermetallic compound FeNi3, both of which exhibit fcc structure, and α-Fe (pearlite). However, the metastable and ordered phases FeNi and Fe3Ni could appear instead. As shown in Figure 3B, Ni strongly stabilised the γ-phase and hence it was assumed that layer 2 represented austenite. This was supported by the appearance of martensite at the steel side of this layer, indicating higher Ms temperature caused by lower Ni content (Figure 3C). Further analyses were proposed to clarify the structure and composition of layer 1. After 15 hours’ sintering two additional layers, 3 and 4, appeared on the hardmetal side of the interlayer, as shown in Figure 6B. EDS and micrographic examinations revealed η-phase in both layers, with increasing Fe content towards the steel side. These findings implied that the occurrence of η-phase was thermodynamically possible in a narrow zone, though no η-phase was formed during conventional sintering. The two-phase field WC + fcc phase of hardmetals with Ni-based binders was located slightly below the stoichiometric carbon composition, while the opposite applied to Fe-based hardmetals. As a result, interdiffusion of Fe led to a strong shift of the twophase field and in turn to carbon deficiency. For this reason Ni-based hardmetal binders were expected to be more prone to the formation of η-phase at the interface than Co-based binders. Sinter-bonded steel/hardmetal composites could be hardened by conventional heat treatment of the appropriate steel without formation of cracks, as the cemented carbides studied had excellent resistance to thermal shock. 100Cr6 showed complete transformation from pearlitic to bainitic microstructure. Figure 7 illustrates the microhardness profile across the interface after hardening, both for 100Cr6-WC/10Co and 100Cr6-WC/8Ni. It was noted that
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MPR February 2010
Figure 6. A: Microstructure at the bonding zone of a 100Cr6-WC/8Ni composite showing two separate interlayers at the interface. The sample was sintered at 1270°C for two hours and etched with V2A solution. B: Same as A but sintered for 15 hours at 1270°C. Two additional interlayers are visible at the hardmetal boundary. C: Magnified view of layer 2 showing martensite needles.
Figure 7. Microhardness profile across the boundary layer of two different steel/hardmetal composites. A: 100Cr6-WC/10Co. B: 100Cr6-WC/8Ni. Both composites were sintered at 1270°C for two hours and subsequently austenitised and tempered. The position “0” was set to the sharp border of WC grains. Indicated HV10 values of the bulk materials represent the average of 10 measurements.
measured values of 100Cr6 remote from the interlayer were consistent with the bulk HV10 values. The relatively large scatter was caused by residual porosity in the steel matrix, though a sharp drop occurred in the binder-rich interlayer close to the hardmetal. In both cases a hardness of only about 200 HV was obtained. The soft, ductile interlayer might have helped in relaxing high and potentially critical internal stresses, enhancing the resilience of the composite. Huber summarised the research as confirming that only high-carbon steels were appropriate for sinter-bonding to hardmetals, as otherwise a brittle phase interlayer was formed. The appearance of a soft Co- or Ni-rich interlayer was explained by the carbon gradient induced by depletion of carbon at the interface. The process, he explained, had several advantages compared to conventional joining techniques: • Cost and time efficiency: sintering of the steel part and joining were combined in a single step, without the need to clean or polish the mating surfaces. Additionally, the use of powder injection moulding (PIM) offered the possibility of near net-
shape manufacturing in high quantities; • Great versatility: complex geometries and a multitude of materials could be processed via PIM; • Mechanical stability: substance-to-substance bonding provided high strength up to elevated temperatures. High stresses at the interface were attenuated by a soft interlayer and tensile stresses in the hardmetal were partly compensated by residual stresses; and • Environmental friendliness: no additives were necessary and no hazardous fumes were generated. With the design limitations mentioned earlier, the sinter bonding technique thus represented an interesting alternative to conventional joining methods, especially for large-scale production and high performance applications. Too long to reproduce here, the paper contains a carefully detailed analysis of the thermodynamics and phase equilibria of reactions and interactions. The interested reader is recommended to study these in the complete transcript, published in the Proceedings of the 17th Plansee Seminar 2009.
metal-powder.net
Contrasting points from Plansee Ken Brookes reports on two very different subjects of considerable practical importance covered at the Plansee hardmetals conference: the sinter bonding of steel and hardmetal and the suppression of WC/Co recrystallisation and grain growth without specific inhibitors...
T
he novel technique for the fabrication of steel/hardmetal composites, presented by Arno Huber of Robert Bosch GmbH’s Corporate Sector Research and Advance Engineering, GerlingenSchillerhöhe, Germany, attracted attention at Plansee. Sinter bonding combines densification of the steel part and its joining to fully sintered hardmetal in a single step. They cannot be sintered simultaneously because the steel would melt at the carbide’s sintering temperature, but substantial interdiffusion at the sintering temperature of the steel achieves a strong atomic bond. Although the basic principle is fairly obvious, there are major problems in commercialising the concept. Most important of these is the differential shrinkage of the two components, causing warpage, high internal stresses or even fracture during the cooling phase from sintering temperature. In brazing, carried out at substantially lower temperatures, the similar effects can be combated by absorption of stress by the braze metal or deformable gauze. In the research discussed in this paper, the carbide component fits loosely (though to a carefully calculated degree) in the steel cavity at sintering temperature, the pressure needed for diffusion bonding being applied by steel shrinkage during cooling. Shrinkage stresses thus augment the sinter bonding. High-carbon steels like 100Cr6 work well in this application. By contrast,
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MPR February 2010
low-carbon steels such as 17-4PH or maraging steels absorb carbon from the carbide constituent of the hardmetal and then form a brittle interlayer consisting of Ȗphase, adversely affecting the mechanical properties of the composite. The microstructural and mechanical characteristics of the bonding zone in several combinations of steels and hardmetals were investigated for this research. Hardmetals, also known as cemented or sintered carbides, exhibit excellent wear resistance, hardness up to around 2000 HV, high transverse rupture strength, high Young`s modulus and compressive strength up to 6000 MPa. However, hardmetals are comparatively brittle, thus prone to breakage caused by shear and tensile stresses, and their material and production costs are high compared to steel. Hence it is common practice to join hardmetal parts to steel supports, typically by brazing, welding, diffusion bonding, overlaying or the use of mechanical methods or adhesives. Some important aspects of different joining methods are listed in Table 1. To study the applicability of sinter bonding to different steel/hardmetal combinations, two different hardmetal grades and four steel grades, listed in Table 2, were selected for the experiments. The hardmetal WC/10Co is a very common grade used in wear applications while the Ni-containing grade has high compression strength and better corrosion resistance,
making it suitable for high-pressure applications. The basic criterion for the steel selection was high yield strength, as the support for hardmetal should be strong and rigid. The steel 100Cr6 exhibits high strength, hardness and wear resistance and is often used for bearings. Where corrosion resistance plays a role, X65Cr13 may substitute for 100Cr6. The high-alloy maraging steel 18Ni350 combines tensile strength of about 2450 MPa with high ductility. Precipitates like Ni3Ti and Ni3Al are responsible. 17-4PH is also a highly alloyed martensitic precipitation-hardening stainless steel, commonly produced by powder injection moulding (PIM) because of its easy sintering and heat treatment. For 100Cr6, two powders with different grain size were used. The finer powder with d50 5.7 μm was used in combination with N09 and the other with OM10. The most important characteristics of the selected steels and hardmetals are listed in Table 2. The sinter bonding experiments employed cylindrical hardmetal inlays and steel green parts with a corresponding cavity. To provide suitable hardmetal inlays, OM10 powder feedstock was prepared for PIM. For this purpose, 59 vol% of hardmetal powder was premixed with an organic binder containing about twothirds polyamide and one-third paraffin wax. This feedstock premix was homogenised and granulated in a twin-screw
0026-0657/10 ©2010 Elsevier Ltd. All rights reserved.
Table 1. Overview of common methods for joining steel and hardmetal. Joining method
Advantages
Disadvantages
Max. shear strength
Brazing
- high reliability - cost and time efficiency
- high residual stresses - insufficient wetting may occur - hazard of toxic fumes caused by solder and fluxing agents
320 MPa
Welding (e.g. electron beam welding, friction welding)
- thermal stresses are minimized compared to brazing - applicable for large parts
- brittle γ-phases, pores and cracks may form in the joint zone
Comparable to brazed joints
Adhesives
- no thermal stresses are induced - bonding zone is thermally and chemically not affected - health hazards and capital costs are low
30 MPa - low strength - only adequate for low operating temperatures
Mechanical attachments (clamping, wedging, screw mounting)
- easy substitution of worn parts - no thermal stresses are generated - retention of optimum heat treatment properties
- no substance-to-substance bond - stress concentrations are induced - geometric limitations
Shrink fit mounting
- homogeneous stress distribution - no substance-to-substance bond - temperature-dependent bond - compressive stresses are strength induced within the hardmetal (pre-load)
Diffusion bonding
- thermal stresses are partly removed by a ductile interlayer - high strength - joining areas may be very large
compounder. Cylindrical bars were injection moulded with a diameter of 6.0mm and cut into 25mm lengths. The green samples were machined to a pencil-like shape with cone angle of 120° and diameter 5.6-5.7mm. Afterwards the samples were cut to a length of about 6mm and the paraffin wax binder mainly removed in acetone (65°C, 1000 min). Thermal debinding was completed and sintering
- mating surfaces must be perfectly planar and clean - time and cost consuming
carried out in a single one-hour hydrogenatmosphere stage up to 1400°C in a Gero graphite furnace. To increase carbon activity in the atmosphere but minimise interdiffusion of carbon, the samples were placed on alumina powder in a closed graphite boat. As Ni-based hardmetals had to be processed in vacuo, sintered inlays of N09 with similar geometry were purchased from Tribo GmbH. In all sintered samples
Depending on geometry and joining method
Depending on oversize
380 MPa
both free graphite and eta phase were absent. The steel feedstocks were prepared in a similar manner, but a powder loading between 59 and 68 vol% was chosen, depending on the powder characteristics. Plate-like specimens were injectionmoulded with a small borehole placed in each centre. A wider, concentric blind hole with a cone angle of 120° was drilled, to serve as a seat for the hardmetal inlay. The
Table 2. Specifications of selected hardmetals and steels. Designation
Composition1 wt%
Grain size2 d50 µm
Strength3 MPa
2.5
not specified
Multi Metals
1
≥ 3,600
Tribo Hartstoff GmbH
5.7 / 8.2
1,600 – 2,400
0.4
Sandvik Osprey Ltd
Elongation %
Supplier
Hardmetals: OM10
WC/10Co
N09
WC/8Ni
Steels: 100Cr6 (1.3505)
Fe-1C-1.6Cr-0.35Mn-0.23Si
X65Cr13 (1.4037) Fe-0.63C-13.8Cr-0.76Mn-0.12N
8.1
570 (cast)
not specified
Sandvik Osprey Ltd
18Ni350
Fe-17.5Ni-12Co-3.3Mo-1.2Ti0.1Mn-0.1Cr-0.08Al
4.2
2,400
6
Sandvik Osprey Ltd
17-4PH (AISI 630)
Fe-17Cr-4.1Cu-4Ni-0.48Mn0.36Si-0.23Nb-0.17Mo-0.11N
8.6
515 – 1,170
10 – 18
Sandvik Osprey Ltd
1 Manufacturer’s data 2 Hardmetals: grain size after sintering; steels: grain size of the powder 3 Hardmetals: transverse-rupture-strength (manufacturer’s data); steels: yield strength σ0.2 of wrought material (literature data)
metal-powder.net
February 2010 MPR
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steel parts were also debound in acetone at 65°C for 1000 minutes. For the sinter bonding process, the hardmetal inlays were inserted into the cavity of the steel green body and the samples sintered in a hydrogen atmosphere. Figure 1 indicates that the gap between steel and hardmetal was closed during sintering, caused by the shrinkage of the steel part, and a strong atomic bond achieved due to interdiffusion. The author explained that the diameter of the cavity had to be chosen so that an apparent oversize of the hardmetal inlay was created during sintering. In this way the surfaces came into close contact, supporting the formation of the bond. The oversize would not occur at the sintering temperature, since diffusional flow or creep would deform the steel part as soon as a contact was established. The apparent oversize could be determined accurately by measuring the total shrinkage of reference samples of the respective steel. The hardmetal inlay had to be able to slide upwards during the sintering process, indicated by the arrows in Figure 1. If sliding were supressed, by too high an inlay weight, for example, bonding at the side surfaces would be poor and the steel part might deform. Thus bigger and heavier parts should be sinter-bonded using a “mounting rack” of some kind. Some of the sintered composites were heat-treated in order to study the influence of interdiffusion between the two components on the microstructure and mechanical properties. The samples were studied with both optical and scanning electron microscopy. Element distribution was analysed by EDS (energy dispersive X-ray spectroscopy). Vickers hardness was measured with a 10kg load and microhardness across the boundary surface, to study the influence of interdiffusion on the mechanical properties of the heat-treated samples. As the hardness values differed heavily the microhardness load was varied between 25g and 500g. As expected, the quality of the steel/ hardmetal bond depended strongly on the apparent oversize of the hardmetal inlay after sintering. Higher oversize resulted in earlier and closer contact between the mating surfaces, which in turn promoted diffusional flow. However, if the oversize were too high, noticeable deformation of the steel part might occur. Optimal oversize depended on the fit between the two parts, their surface finish and the materials. Oversize values between 1% and 5%
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MPR February 2010
Figure 1. Schematic illustration of the sinter bonding process. The sintered hardmetal inlay is inserted into the cavity of the steel support green body. The gap between both parts is closed upon sintering and a strong bond develops due to interdiffusion.
seemed reasonable. Diffusion mechanisms played a crucial role in bonding and oxide layers could significantly impede interdiffusion. Fortunately, oxides could generally be removed at an early stage of sintering by applying a reducing hydrogen atmosphere or vacuum, avoiding costly cleaning of the mating surfaces. Due to interdiffusion between the steel and hardmetal binder phase, a solid solution occurs first at the interface and then intermetallic compounds or precipitates depending on the compositions of both materials. In this regard it was advantageous that Fe, Co and Ni exhibited face-centred cubic (fcc) structures and were completely intersoluble at sintering temperature. The sintered 100Cr6-WC/10Co composite samples showed excellent bonding between steel and hardmetal inlay, illustrated in Figure 2A. A narrow porous layer formed near the interface, indicating strong interdiffusion between the parts. Adjacent to the interface, a bright, 300 – 400μm thick, zone was visible in the hardmetal component. During cooling, high residual stresses were generated in the composite due to the significant difference between the coefficients of thermal expansion (CTE) of steel and hardmetal. The greater contraction of steel led to high compressive stresses in the hardmetal inlay which was beneficial for the durability of the assembly. First, the
compressive stresses generated a clamping force which counteracted potential failure of the bonding zone. Second, tensile stresses in the hardmetal inlay, which could be crucial and might have arisen from the high internal pressure, were partly compensated. However, if the composite was cut axially the stress state was altered and cracks formed in the hardmetal inlay, as shown in Figure 2B. The cracks did not proceed along the boundary but within the hardmetal part, demonstrating high strength in the bond. 100Cr6 had a fully pearlitic microstructure and a bright and diffuse interlayer between both parts, shown in Figure 2C. EDS revealed that this layer predominantly contained cobalt and iron (Figure 3A). The penetration depth of Co and Fe into the steel and hardmetal sides respectively was about 70 – 80μm. In this line scan no diffusion of W could be observed, but the comparison of spectra of different EDS maps showed that some W was present in the Co-rich interlayer. Figure 3B shows a simplified phase diagram of the Fe/Co/Ni system. Co and Ni both stabilise Ȗ-Fe (gamma iron) with fcc structure and austenite is retained at room temperature if these elements are present in sufficiently high amounts. Martensite is formed only in a narrow range near the Ȗ-phase boundary upon cooling. W, Cr and C, also present within the interlayer,
Figure 2. Optical micrographs of the microstructure at the boundary between 100Cr6 and WC/10Co after sintering. The samples were sintered at 1270°C for two hours under hydrogen atmosphere and 100Cr6 showed relative density of 95.8%. A: horizontal section. B: axial section. C: microstructure etched with V2A-solution (room temperature, 180s).
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decrease the martensite transformation temperature Ms and also the martensite content. This is in agreement with microscopic examinations as no martensite was observed. However, no distinct phase boundary was visible at the interface as would be expected at a pearlite/γ-phase boundary. Instead a smooth transition of pearlitic to a more homogeneous structure was observed, showing small precipitates containing tungsten. This implied that the interlayer had a body-centred cubic (bcc) structure.
Cobalt enrichment at the interface was not expected, as concentration or activity gradients should have been eliminated by diffusion. To study this effect, hardmetal inlays were heat-treated under the same sintering conditions as the composites, avoiding the influence of interdiffusion between steel and hardmetal. After sintering, a 10 – 30μm thick surface layer had formed, most pronounced at the edges, indicating the influence of the sintering atmosphere (Figure 4A).
Figure 3. A: EDS analysis of the element distribution across the interface between 100Cr6 and WC/10Co (1270°C, 2 h). The interface line indicates the borderline of WC grains. In order to reduce scattering each data point was derived from an EDS analysis map by averaging a 35 μm broad zone perpendicular to the line axis. B: Fe-Co-Ni system showing composition range and temperature of martensite transformation.
Hardmetals tend to form so called Ȗphases (M6C or M12C, generally double carbides of the hard constituent and the binder alloy) if the carbon content falls slightly below the stoichiometric value, illustrated in Figure 4B. These phases, brittle and therefore undesirable, can be seen after etching with Murakami solution (equal parts of NaOH and K4[Fe(CN)6]). In Figure 4A the surface layer contains both Ȗphases and a second phase identified as Co binder. It appeared that the carbon-poor atmosphere was responsible for the formation of a Co-rich surface layer, not observed under common (carbon-rich) sintering conditions for hardmetals. Diffusion of cobalt into regions of lower carbon content in hardmetals had been verified by several previous studies. Hence it was assumed that the combined interaction with the atmosphere and the steel was responsible for Co enrichment at the interface. Sintered composite samples were etched to show Ȗphases at the interface, but at most only a minor fraction of dispersed Ȗphase was observed, both in the Co-rich interlayer and in the hardmetal close to the interface. This was surprising, as carbon diffuses interstitially and was expected to be orders of magnitude faster than substitutionally diffusing elements, leading to high diffusional exchange as soon as contact between steel and hardmetal was established.
Higher carbon content
Figure 4. A: Hardmetal inlay showing a surface layer which consists of Co binder phase (bright) and γ-phase (dark). The sintered sample was tempered at 1270°C for two hours under a hydrogen atmosphere, polished and etched with Murakami solution at room temperature for 10 seconds. B: Isothermal section of the W-C-Co phase diagram at 1350°C.
Figure 5. Microstructure at the bonding zone of different steel/hardmetal-composites showing a brittle interlayer. A: Optical micrograph of X65Cr13-WC/10Co (1270°C, two hours). B: SEM micrograph of 18Ni350-WC/10Co (1350°C, one hour). C: SEM-micrograph of 18Ni350 near the interlayer with different precipitates at the grain boundaries, indicated by the arrows (1350°C, one hour).
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In contrast to 100Cr6, X65Cr13 and the nearly carbon-free steels 18Ni350 and 17-4PH showed a brittle interlayer between steel and hardmetal, seen in Figure 5. This layer mainly contained Co, Fe, W and C as the brittle Ȗphase. Because X65Cr13 contained substantially higher carbon content than the other two steels, interdiffusion of carbon was less pronounced, resulting in a thinner interlayer. Small cracks were visible in the interlayer, due to high residual stresses and precipitates appeared on the steel side near the interface because of interdiffusion. 18Ni350 in particular showed precipitates at the grain boundaries (Figure 5C), which acted as nucleation sites or as preferred diffusion paths for carbon. The bright precipitates contained mainly Mo, W and C and the darker ones Ti and C. It was assumed that these precipitates were carbides. In addition to WC/10Co, a hardmetal with a nickel binder was studied in combination with 100Cr6. A layer consisting of
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Ni binder and Ȗphase was observed on the free surface of the hardmetal inlays after sintering. A relatively broad interdiffusion zone formed between steel and hardmetal, as shown in Figure 6A, subdivided into two separate layers 1 and 2. As with the Co-based hardmetal, a binder-rich interlayer was formed, due to the carbon gradient from the carbon-deficient interface. At the sintering temperature of 1270°C, iron and nickel form a solid solution with fcc structure over the whole concentration range. However, at room temperature several phases are thermodynamically stable: γ-Ni and the intermetallic compound FeNi3, both of which exhibit fcc structure, and α-Fe (pearlite). However, the metastable and ordered phases FeNi and Fe3Ni could appear instead. As shown in Figure 3B, Ni strongly stabilised the γ-phase and hence it was assumed that layer 2 represented austenite. This was supported by the appearance of martensite at the steel side of this layer, indicating higher Ms temperature caused by lower Ni content (Figure 3C). Further analyses were proposed to clarify the structure and composition of layer 1. After 15 hours’ sintering two additional layers, 3 and 4, appeared on the hardmetal side of the interlayer, as shown in Figure 6B. EDS and micrographic examinations revealed η-phase in both layers, with increasing Fe content towards the steel side. These findings implied that the occurrence of η-phase was thermodynamically possible in a narrow zone, though no η-phase was formed during conventional sintering. The two-phase field WC + fcc phase of hardmetals with Ni-based binders was located slightly below the stoichiometric carbon composition, while the opposite applied to Fe-based hardmetals. As a result, interdiffusion of Fe led to a strong shift of the twophase field and in turn to carbon deficiency. For this reason Ni-based hardmetal binders were expected to be more prone to the formation of η-phase at the interface than Co-based binders. Sinter-bonded steel/hardmetal composites could be hardened by conventional heat treatment of the appropriate steel without formation of cracks, as the cemented carbides studied had excellent resistance to thermal shock. 100Cr6 showed complete transformation from pearlitic to bainitic microstructure. Figure 7 illustrates the microhardness profile across the interface after hardening, both for 100Cr6-WC/10Co and 100Cr6-WC/8Ni. It was noted that
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MPR February 2010
Figure 6. A: Microstructure at the bonding zone of a 100Cr6-WC/8Ni composite showing two separate interlayers at the interface. The sample was sintered at 1270°C for two hours and etched with V2A solution. B: Same as A but sintered for 15 hours at 1270°C. Two additional interlayers are visible at the hardmetal boundary. C: Magnified view of layer 2 showing martensite needles.
Figure 7. Microhardness profile across the boundary layer of two different steel/hardmetal composites. A: 100Cr6-WC/10Co. B: 100Cr6-WC/8Ni. Both composites were sintered at 1270°C for two hours and subsequently austenitised and tempered. The position “0” was set to the sharp border of WC grains. Indicated HV10 values of the bulk materials represent the average of 10 measurements.
measured values of 100Cr6 remote from the interlayer were consistent with the bulk HV10 values. The relatively large scatter was caused by residual porosity in the steel matrix, though a sharp drop occurred in the binder-rich interlayer close to the hardmetal. In both cases a hardness of only about 200 HV was obtained. The soft, ductile interlayer might have helped in relaxing high and potentially critical internal stresses, enhancing the resilience of the composite. Huber summarised the research as confirming that only high-carbon steels were appropriate for sinter-bonding to hardmetals, as otherwise a brittle phase interlayer was formed. The appearance of a soft Co- or Ni-rich interlayer was explained by the carbon gradient induced by depletion of carbon at the interface. The process, he explained, had several advantages compared to conventional joining techniques: • Cost and time efficiency: sintering of the steel part and joining were combined in a single step, without the need to clean or polish the mating surfaces. Additionally, the use of powder injection moulding (PIM) offered the possibility of near net-
shape manufacturing in high quantities; • Great versatility: complex geometries and a multitude of materials could be processed via PIM; • Mechanical stability: substance-to-substance bonding provided high strength up to elevated temperatures. High stresses at the interface were attenuated by a soft interlayer and tensile stresses in the hardmetal were partly compensated by residual stresses; and • Environmental friendliness: no additives were necessary and no hazardous fumes were generated. With the design limitations mentioned earlier, the sinter bonding technique thus represented an interesting alternative to conventional joining methods, especially for large-scale production and high performance applications. Too long to reproduce here, the paper contains a carefully detailed analysis of the thermodynamics and phase equilibria of reactions and interactions. The interested reader is recommended to study these in the complete transcript, published in the Proceedings of the 17th Plansee Seminar 2009.
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