Technical Field

The present invention relates generally to the field of metallurgy and bearing steels. The bearing steel exhibits improved resistance to hydrogen embrittlement.

Background

Bearings are devices that permit constrained relative motion between two parts. Rolling element bearings comprise inner and outer raceways and a plurality of rolling elements (balls or rollers) disposed therebetween. For long-term reliability and performance it is important that the various elements have a high resistance to rolling contact fatigue, wear and creep.

Steelmaking companies have been active in lowering the hydrogen content during casting, since this element can have an adverse effect on the rolling contact fatigue life. The hydrogen concentration should typically not exceed 1 ppm. Even if the hydrogen content is very low in the as-produced steel, its amount is likely to increase during service, for example due to oil decomposition or electric current breaking through the layer of oil, resulting in the decomposition of oil molecules into products including free hydrogen, making its ingress into the bulk possible. Hydrogen embrittlement is likely to occur when the steel contains mobile hydrogen. For this reason it has been proposed to immobilise hydrogen in the alloy microstructure.

The steel known as 10006 has the following composition: 0.974 wt% carbon, 0.282 wt% silicon, 0.276 wt% manganese, 0.056 wt% molybdenum, 1.384 wt% chromium, 0.184 wt% nickel, 0.042 wt% aluminium, 0.21 wt% copper, 0.01 wt% phosphorus and 0.017 wt% sulphur, the balance being iron (and any unavoidable impurities). This steel exhibits high hardness and is suitable for use in a bearing component. However, 100Cr6 exhibits moderate-to-low resistance to hydrogen embrittlement. Bainitic steel alloys are produced by the transformation of austenite to ferrite at intermediate temperatures of typically from 200 to 550 ^° C. The cooling of the austenite leads to a microstructure comprising ferrite, carbides and retained austenite. Bainite itself comprises a structure of supersaturated ferrite containing particles of carbide, the dispersion of the latter depending on the formation temperature. The hardness of bainite is usually somewhere intermediate between that of pearlite and martensite.

Superbainite alloys have been produced without the presence of carbides (or at least with a very small volume fraction of carbides). That is, the structure is formed of a fine mixture of bainitic ferrite and carbon-enriched retained austenite. This produces a combination of high toughness (typically up to 40 MPam ¹ ' ² ) and high tensile strength (typically around 2500 MPa).

Superbainite alloys are produced by austenitisation followed by prolonged isothermal transformation at a temperature slightly above the martensite start (Ms) temperature.

Lower transformation temperatures are desirable because they allow for the formation of a finer structure having improved mechanical properties. However, lower temperatures result in longer processing times. It is an object of the present invention to address or at least mitigate some of the problems associated with prior art, or at least to provide a commercially useful alternative thereto.

Summary

The present invention provides for the use of a steel alloy having a microstructure comprising superbainite in a bearing component exposed to a hydrogen-containing environment to achieve a reduction in hydrogen embrittlement. Hydrogen-containing environments include any hydrogen-containing atmosphere, either gaseous or liquid. Examples are oils and/or greases.

The bearing component is formed from a steel alloy composition which preferably comprises: from 0.6 to 1.0 wt% carbon,

from 1 .0 to 2.0 wt% silicon,

from 1.5 to 2.5 wt% manganese,

from 0.1 to 0.5 wt% molybdenum,

from 0.5 to 1 .5 wt% chromium,

from 0.5 to 1.5 wt% aluminium,

from 1 .0 to 2.0 wt% cobalt, optionally one or more of the following elements from 0 to 0.5 wt.% nickel,

from 0 to 0.2 wt.% vanadium,

from 0 to 0.3 wt.% copper,

from 0 to 0.1 wt.% sulphur,

from 0 to 0.1 wt.% phosphorous, and the balance iron, together with any unavoidable impurities.

The reduction in hydrogen embrittlement is related to an increase in hydrogen trapping capacity of the alloys according to the present invention.

A typical superbainitic structure comprises untransformed retained austenite and bainitic ferrite. Typically, the main structural feature is the fineness of the bainitic ferrite platelets, which are typically tens of nanometres thick. The alloy preferably comprises bainite as the main phase (typically at least 60% bainite, more typically at least 80% bainite) or as essentially the only phase (i.e. > 95% bainite). Bainite is preferably obtained by carrying out the transformation at a relatively low temperature. One result of the low transformation temperature is that, as noted above, the plates of bainite are very fine. In particular, the material preferably has a microstructure comprising plates of bainite of less than 100 nm, typically from 10 to 50 nm, more typically from 20 to 40 nm.

The plates of bainite are typically interspersed with retained austenite. The bainite typically forms at least 60% of the microstructure, more typically at least 80%. The bainite is preferably lower bainite. The steel is preferably essentially carbide-free. Typically, the microstructure will comprises less than 5% carbides, more typically less than 3%.

The steel typically has an ultimate tensile strength of 2500 MPa, a hardness at 600-670 HV, and toughness in excess of 30-40 MPam ¹ ' ² . The microstructure and resulting mechanical properties lead to improved rolling contact fatigue performance in the bearing component.

The present invention will now be further described. In the following passages different aspects of the invention are defined in more detail. Each aspect so defined may be combined with any other aspect or aspects unless clearly indicated to the contrary. In particular, any feature indicated as being preferred or advantageous may be combined with any other feature or features indicated as being preferred or advantageous. The steel composition preferably comprises 0.5 to 1 .2 wt% carbon, more preferably 0.6 to 1.1 wt% carbon. When the carbon content is higher than about 1.1 wt%, there is a reduction in the maximum volume fraction of the bainitic ferrite portion of the

microstructure. When the carbon content is lower than about 0.6 wt%, the alloys have a higher martensite start temperature (Ms). The martensite start temperature acts as a lower limit for conducting the superbainite isothermal transformation and thus constrains the low temperature superbainite formation. Preferably, the steel composition comprises from 0.7 to 1 wt% carbon, more preferably 0.7 to 0.9. Most preferably, the steel composition comprises about 0.8 wt% carbon. The steel composition preferably comprises 0.15 to 2 wt% silicon, more preferably 0.25 to 2 wt% silicon, still more preferably 1 to 2 wt% silicon. In combination with the other alloying elements, this results in the desired fine carbide-free microstructure (or essentially carbide-free). The addition of silicon is advantageous because it suppresses the formation of carbides (cementite). If the silicon content is lower than about 1 wt%, then cementite may be formed at low temperatures preventing the formation of superbainite. However, too high a silicon content (for example above about 2 wt%) may result in undesirable surface oxides and a poor surface finish. Preferably, the steel composition comprises 1 .3 to 1 .7 wt% silicon and, more preferably, 1.4 to 1 .6 wt% silicon. The steel composition preferably comprises 0.25 to 2.5 wt% manganese. Manganese acts to increase the stability of austenite relative to ferrite. Manganese also acts to increase the hardenability of the alloy. Preferably, the steel composition comprises 1 .5 to 2.5 wt% manganese and, more preferably, 1.8 to 2.2 wt% manganese, still more preferably 1.9 to 2.1 wt% manganese.

The steel composition preferably comprises 0.5 to 3 wt% chromium, preferably 0.5 to 2 wt% chromium, more preferably 0.6 to 1 .5 wt% chromium, still more preferably 0.8 to 1.2 wt% chromium. Chromium acts to increase hardenability and reduce the bainite start temperature. Chromium also has a corrosion-resistant effect.

While cobalt and aluminum are optional elements, it is preferable for one or both elements to be present. Accordingly, in a preferred embodiment, the steel composition preferably comprises one or both of: from 0.1 to 5 wt% cobalt, and/or

from 0.1 to 2 wt% aluminium. More preferably, the steel composition comprises one or both of: from 1 to 3 wt% cobalt, and/or

from 0.5 to 2 wt% aluminium. More preferably, the steel composition comprises one or both of: from 1 to 2 wt% cobalt, and/or

from 0.5 to 1.5 wt% aluminium. Still more preferably, the steel composition comprises one or both of: from 1.3 to 1 .7 wt% cobalt, and/or

from 0.8 to 1.2 wt% aluminium. Aluminium has been found to improve the intrinsic toughness of the bearing component, possibly due to it suppressing carbide formation. However, the use of aluminium requires stringent steel production controls to ensure cleanliness and this increases the processing costs. In one embodiment, the steel composition comprises substantially no aluminium.

Cobalt has been found to improve the corrosion resistance of the bearing component. This is very important for bearing components for wind turbines or marine pods, for example. Such bearings may become contaminated by sea water, which can drastically reduce the service life of the bearing. The addition of cobalt further increases the rate of superbainite formation.

If present, the alloy preferably also comprises 0.05 to 0.5 wt% molybdenum.

Molybdenum acts to avoid austenite grain boundary embrittlement owing to impurities such as, for example, phosphorus. Molybdenum also acts to increase hardenability and reduce the bainite start temperature. The steel composition preferably comprises 0.1 to 0.5 wt% molybdenum, more preferably 0.1 to 0.4 wt% molybdenum. The molybdenum content in the alloy is preferably no more than about 0.3 wt% otherwise the austenite transformation into bainitic ferrite may cease too early, which can result in significant amounts of austenite being retained in the structure. Preferably, the steel composition comprises 0.15 to 0.3 wt% molybdenum and, more preferably, 0.2 to 0.3 wt%

molybdenum.

Silicon and aluminium additions both act to suppress carbide formation. The

superbainite formation is encouraged by the suppression of cementite. Accordingly, in one embodiment, the total amount of aluminium and silicon is 1 to 2.5 wt%, more preferably 1.5 to 2 wt%.

The steel composition optionally comprises up to 0.25 wt% vanadium. In one embodiment, the steel composition comprises from 0.1 to 0.2 wt% vanadium. Vanadium can be useful during austenitisation because it helps to control the austenite grain size. In another embodiment, the steel composition comprises substantially no vanadium.

A preferred steel alloy according to the present invention comprises: from 0.6 to 1.0 wt% carbon

from 1 .0 to 2.0 wt% silicon

from 1.5 to 2.5 wt% manganese

from 0.1 to 0.5 wt% molybdenum

from 0.5 to 1 .5 wt% chromium

from 0.5 to 1.5 wt% aluminium

from 1 .0 to 2.0 wt% cobalt optionally one or more of the following elements from 0 to 0.5 wt% nickel

from 0 to 0.2 wt% vanadium

from 0 to 0.3 wt% copper

from 0 to 0.1 wt% sulphur

from 0 to 0.1 wt% phosphorous and the balance iron, together with unavoidable impurities. More preferably, the steel alloy according to the present invention comprises: from 0.7 to 0.9 wt% carbon

from 1 .3 to 1.7 wt% silicon

from 1.7 to 2.3 wt% manganese

from 0.1 to 0.4 wt% molybdenum

from 0.8 to 1 .2 wt% chromium

from 0.8 to 1.2 wt% aluminium

from 1 .3 to 1.7 wt% cobalt optionally one or more of the following elements from 0 to 0.5 wt% nickel

from 0 to 0.2 wt% vanadium

from 0 to 0.3 wt% copper from 0 to 0.1 wt% sulphur

from 0 to 0.1 wt% phosphorous and the balance iron, together with unavoidable impurities.

As noted above, the steel alloy has a superbainitic microstructure.

It will be appreciated that the steel alloys according to the present invention may contain unavoidable impurities, although, in total, these are unlikely to exceed 0.5 wt% of the composition. Preferably, the alloys contain unavoidable impurities in an amount of not more than 0.3 wt% of the composition, more preferably not more than 0.1 wt% of the composition. Phosphorous and sulphur are preferably kept to a minimum.

The steel alloys according to the present invention may consist essentially of the recited elements. It will therefore be appreciated that in addition to those elements that are mandatory other non-specified elements may be present in the composition provided that the essential characteristics of the composition are not materially affected by their presence.

The steel composition of the present invention enables the formation of superbainite. This structure typically has a fine-scale bainitic-ferrite with excellent mechanical properties. It is achieved by transforming austenite to bainite at a relatively low temperature, typically from 150 to 350 ^° C.

The present inventors have discovered an advantageous blend of alloying elements that allows for the formation of a superbainitic steel composition at lower temperatures. As a result, when transformed at low temperatures, the mechanical properties of the product are substantially improved. Alternatively, products of equivalent properties to known superbainite materials may be produced with significantly reduced processing times.

Preferably at least 50 vol% of the microstructure comprises bainitic ferrite, more preferably at least 70 vol%, still more preferably at least 90 vol%. In one embodiment, essentially all of the structure is bainitic ferrite. That is at least 95 vol%, more preferably at least 98 vol%, more preferably at least 99 vol% of the structure comprises bainitic ferrite. The bainitic ferrite typically has a very fine structure.

The steel composition of the present invention preferably has a structure comprising at most 10 vol% carbides, more preferably at most 5 vol%, still more preferably at most 2 vol%. In a preferred embodiment, the structure is substantially carbide-free.

The superbainite structure is easily characterised using high-resolution techniques. It comprises a fine mixture of bainitic ferrite and carbon-enriched retained austenite. In one aspect, the average size of the sandwiched platelets of ferrite/austenite is typically less than 50 nm, for example 20 to 40 nm, more typically 25 to 35 nm.

In an embodiment, up to 50 vol% of the microstructure comprises retained austenite. More particularly, 1 to 50 vol% of the structure comprises retained austenite, more preferably 5 to 45 vol%, still more preferably 5 to 40 vol%, still more preferably 5 to 30 vol%. The retained austenite is typically carbon-enriched.

In another embodiment, the microstructure advantageously comprises 5 to 20 vol% retained austenite, optionally up to 5 vol% carbides, and the balance comprising bainitic ferrite. More preferably, the structure comprises 10 to 20 vol% retained austenite, optionally up to 2 vol% carbides, and the balance comprising bainitic ferrite.

The superbainitic microstructure has been found to be advantageous in that it increases hydrogen trapping capacity and thus reduces hydrogen embrittlement.

Rolling element bearings comprise inner and outer raceways and a plurality of rolling elements (balls or rollers) disposed there-between. For long-term reliability and performance it is important that the various elements have resistance to rolling fatigue, wear and creep. In the present invention, a bearing component is formed from a steel alloy as herein described. Examples of bearing components where the steel can be used include a rolling element (ball or cylinder), an inner ring, and an outer ring. The present invention also provides a bearing comprising a bearing component as herein described. According to a further aspect, there is provided a method for preparing a steel alloy as herein described, the method comprising:

(i) providing a steel composition as herein described;

(ii) heating the steel composition to at least partially austenitise the composition; and

(iii) maintaining the composition at a superbainite formation temperature until a superbainite structure is formed.

The method may further comprise following step (iii):

(iv) heating the composition at a temperature greater that the superbainite formation temperature to reduce the amount of retained austenite. Step (iv), which is optional, may be conducted at a temperature of typically 450 to 550 ^° C.

The step of heating the composition to at least partially austenitise the composition will typically involve raising the temperature of the composition to at least about 880 ^° C, preferably from 880 to 960 ^° C, more preferably from 900 to 940°C. The step of maintaining the composition at a superbainite formation temperature until a superbainite structure is formed will typically comprise adjusting the temperature of the hot austenitised composition to a temperature of from 150 to 400 ^° C. The preferred heat- treatment is from 150 to 350 ^° C. The treatment is preferably, but not necessarily, isothermal. This leads to the most consistent structure. The processing time is typically 1 to 100 hours, preferably 6 to 72 hours, although longer times are possible. As far as resistance to hydrogen embrittlement is concerned, it has been found that a temperature of at least 200°C, and preferably at least 300°C, results in an improvement in hydrogen trapping capacity. The structure of the steel alloy described herein can be determined by conventional microstructural characterisation techniques such as, for example, optical microscopy, TEM, SEM, AP-FIM, and X-ray diffraction, including combinations of two or more of these techniques. Figures

The present invention will now be described further in relation to the figures annexed hereto by way of non-limiting examples.

Figure 1 shows a possible processing route and heat-treatment for a steel alloy used in the present invention.

Figure 2 shows the heat treatment schedules of a) SB200 (invention), b) SB300

(invention), and c) 10006 (comparison).

Figure 3 shows melt extraction results showing the a) total hydrogen measured just after H-charging, and b) trapped hydrogen content determined 24 hours after H-charging. Figure 4 shows melt extraction results showing hydrogen content in superbainite transformed at 300°C and 200°C in Figure 2. The experiment was performed after H- charging.

According to the schematic shown in Figure 1 , in a first step 1 an alloy composition is prepared and cast. The alloy composition is then subjected to a conventional high temperature soaking step 2, which is followed by a hot-rolling step 3.

The hot rolling step 3 is typically carried out at a starting temperature of 1 150°C. Several hot-rolling passes may be applied as necessary. The hot-rolled steel, which can be in a bar or plate form, is then allowed to cool slowly to room temperature to avoid the formation of high-carbon martensite. A typical preferred structure in the as hot-rolled condition, at room temperature, is pearlite.

The hot-rolled material may then optionally be homogenised in a homogenisation step 4, such as a 1200°C treatment for 24 to 48 hours in a vacuum. Optionally, the bars may then be furnace cooled in a cooling step 5, to allow them to cool down slowly to room temperature, also under vacuum. If desired, the material can then be machined in a machining step to near-net-shape components, for example bearing components (this optional step is not shown). The material or machined shapes are subsequently fully or partially austenitised. For example, a 900°C austenitising heat-treatment for 30 minutes (step 6). The

austenitisation may be full or partial, depending on the desired carbon content in the austenite phase.

Immediately after austenitisation, the material (or component) is cooled down to the bainite transformation regime and allowed to transform isothermally (step 7).

Superbainite transformation can be carried out at a fixed transformation temperature, for example at 200 to 300°C for, for example, 6 to 72 hours. A multi-step transformation temperature schedule may also be adopted to tailor the phase fractions in the structure.

After the isothermal treatment step 7, the material (or component) is air-cooled (step 8) and then finished by machining and/or grinding to the required final dimensions (step 9). Finally, an optional step of honing, lapping or polishing can be performed (step 10).

Examples

The invention will now be explained with reference to the following non-limiting

examples.

Composition according to the present invention (SB200 and SB300; SB200A and

SB300A) and comparative compositions (Iron and 100Cr6) were prepared as described in Table 1 below (all wt%).

Table 1 (all wt%)

Alloy C Si Mn Mo Cr Ni Al Cu Co Fe

Iron 0.002 . . . . 0.02 - - balance

100Cr6 0.97 0.28 0.28 0.056 1.38 0.18 0.042 0.21 - balance

SB200 0.82 1 .56 2.03 0.25 0.98 - 1.01 1 .52 balance SB300

SB200 ^A 0.82 1 .56 2.03 0.25 0.98 - 1.01 - 1 .52 balance SB300 ^A

With regard to phosphorus and sulphur in the alloys: for iron, P < 0.005 and S < 0.004; for 100Cr6, P < 0.01 and S < 0.017; and for SB200, SB300, SB200 ^A and SB300 ^A , P < 0.002 and S < 0.001 .

SB200 and SB300 steel alloys were austempered and isothermally transformed at 200°C and 300°C, respectively, leading to an essentially carbide-free, bainitic nanostructure.

100Cr6 was austenitized, quenched to form martensite, and tempered. The chemical compositions of the tested grades in Table 1 was determined using a glow discharge atomic emission spectrometer, LECO GDS850A.

The (pure) iron samples were swagged to 3 mm diameter rods and cut into 10 mm length cylinders. Swaging was chosen because of the initial sample geometry was irregular and difficult to machine. After swaging, the microstructure was highly deformed, so in order to reduce the number of dislocations and increase grain size, the specimens were treated at 900°C for 5 min.

SB200 and SB300 were treated to the schedule presented in Figure 2a and 2b,

respectively. The heat treatments were conducted in two air furnaces, the first kept at 920°C and the second at 200°C or 300°C. SB200 ^A and SB300 ^A were subjected to the same heat treatment as SB200 and SB300, but they were additionally heated up to

500°C and kept for 1 hour (SB300 ^A ) or 2 hours (SB200 ^A ), and then cooled down to room temperature for reducing the retained austenite amount. After heat treatment, the samples were cut into rods of 4 mm diameter and 12 mm length.

The 100O6 grade was spheroidized, and then hardened according to the schedules shown in Figure 2c. The sample was cut into 4 mm diameter and 12 mm length rods, and heat-treated in vacuum on a Thermecmaster dilatometer with helium quenching gas at a cooling rate of 25°Cs ^"1 . The specimens were subsequently cut into 6 mm length rods for hydrogen charging, melt extraction (MET) and thermal desorption analysis (TDA) experiments.

With regard to hydrogen charging of the samples, a wire was spot welded to the specimen which was placed into a charging cell filled with electrolyte (1 dm ³ distilled H ₂ 0, 4 g NaOH, 4 g Thiourea) and surrounded by a platinum wire. The polarity of the sample was negative. Subsequently, the cell was connected to an 8 mA current source. The charging process underwent 24 hours during which the electrolyte was stirred and kept at 80°C. After charging, the samples were gently polished and ultrasonically cleaned with petroleum ether and acetone, respectively, thereby removing the dark layer which appears after charging.

An Eltra ONH-200 hydrogen content analyzer was employed for melt extraction technique (MET). The specimens (approx. 0.5 g weight) were tested immediately and 24 hours after hydrogen charging.

The samples were heated up to their melting point temperature for evaporating all gases, which were then taken via a nitrogen carrier gas through chemical reaction tubes to analyzing units, which measure the thermal conductivity of the carrier gas. The hydrogen content was measured first on a reference sample containing about 6 ppm of hydrogen, which was employed for calibration, and subsequently the experimental samples were measured.

A Pyroprobe 5000 Series, CDS Analytical Inc unit, heated up the specimens at a rate of 2.6°C/min, and hydrogen content was detected by a pulsed discharge detector with helium carrier gas in an Agilent Technologies 7890A GC system with an Agilent 19091 P- MS7 15 m x 320 μηη x 25 μηη (length x external x internal diameter) column used for hydrogen, oxygen and nitrogen gases separation. The samples were analyzed in sequential 3 minutes intervals.

The amount of retained austenite in the examined specimens was measured by X-ray diffraction. 8 mm diameter and 3 mm long samples were ground using 1200 grit SiC paper for obtaining a smooth surface. The grinding process was followed by etching with concentrated HCI for 20 seconds for removing surface deformation caused by grinding. A Bruker AXS BV D8 X-ray diffractometer was employed. A standard DQUANT software was used to quantify the retained austenite amount. X-ray spectra were obtained between 35° and 105° 2Θ with a step size of 0.02° and a dwell time of 10 s. A Mo target was used (MoKa). The determination of retained austenite in martensite is based on the ratio method for the intensities of four peaks, martensite (200) and (21 1 ), and austenite (220) and (31 1 ).

MET analysis performed immediately after hydrogen charging shows the total hydrogen content absorbed by the samples, including diffusible hydrogen (Figure 3a). The analysis conducted 24 hours after charging highlights the hydrogen trapped, with the exception of diffusible hydrogen (Figure 3b), as diffusible hydrogen is thought to have already escaped. SB300 and SB200 showed a much smaller gap between the two conditions compared to100Cr6. SB300 showed no difference. MET on pure iron was also performed for comparison, since ferrite in pure iron resembles bainitic ferrite in superbainite, and may be employed as a reference.

The trapping character of superbainite was subsequently examined on specimens with a lower retained austenite content 24 hours after charging (Figure 4). Retained austenite was decreased by an additional annealing step at 500°C for 1 hour (SB300 ^A ) and 2 hours (SB200 ^A ). The retained austenite amount decreased from 19% to 16% for SB300 and from 23% to 14% for SB200. Longer annealing resulted in lower retained austenite content.

SB200 ⁰ and SB300 ⁰ show the hydrogen content present in the steels that were not subjected to charging before MET analysis.

Thermal desorption analysis was employed to determine the nature of the prevailing hydrogen trap active at a given temperature. 100Cr6, SB300 and SB200 showed different T _TDA peaks. T _TDA =188°C in 100O6 is ascribed to trapping by dislocations. For the martensitic steels, the thermal desorption peaks did not change between the conditions just after charging and 24 hours after charging. In contrast, T _TDA for SB300 changed from 188 to 227°C, and for SB200 from 188 to 219°C, immediately after charging, and 24 hours after it. Trapping in superbainite can be associated with the retained austenite/ferrite interface. T _T DA peaks shift both in SB300 and SB200 after charging, suggesting diffusible or weakly trapped hydrogen disappears with time. The area under the thermal desporption curves (not shown) decreased for all steels 24 hours after charging. For SB300, it decreased from 2.14 to 1 .64 ppm (24%). It is believed that almost all hydrogen is trapped effectively by retained austenite/ferrite interfaces in SB300 immediately after entering the specimens. For the other steel grades, the area under the thermal desporption curves (not shown) decreased as follows: from 2.21 to 0.74 ppm (67% reduction) for SB200; and from 1 .19 to 0.12 ppm (90% reduction) for 100Cr6. The heights of the peaks provide an indication of the effectiveness of the microstructure for hydrogen trapping capacity: SB200 and SB300 exhibit significantly higher peaks than 100Cr6.

SB300 and SB200 showed a lower difference in the amount of absorbed and trapped hydrogen, resulting in a lower content of potentially harmful diffusible hydrogen (see Figure 3). Superbainite showed the lowest degree of hydrogen absorption (see Figure 4).

Since superbainite combines nanostructured ferrite and retained austenite, both phases were examined for their response to hydrogen. Pure iron after charging showed the lowest hydrogen absorption when it was subjected to the same charging conditions as the other tested steels (see Figure 3), suggesting that hydrogen could diffuse throughout the pure iron sample, that it was not immobilised by its microstructure, and that it diffused out during the charging process itself. The pure iron sample appeared to be nearly "transparent" to hydrogen. The influence of retained austenite on hydrogen absorption was tested in SB300 and

SB200 with a decreased amount of such phase. This was achieved by an additional annealing step for retained austenite decomposition at 500°C. The obtained results showed that the overtempered specimens (SB300 ^A and SB200 ^A ) absorbed less hydrogen compared to regular SB300 and SB200 (see Figure 4). This observation suggests that it is the retained austenite-ferrite interface that traps hydrogen in superbainite, other than ferrite or retained austenite themselves. The nature of the retained austenite interface is thought to be important, as it determines whether hydrogen is trapped reversibly or irreversibly. The T _TDA peaks for SB300 and SB200 at 227°C and 219°C, 24 hours after charging, indicate that hydrogen is trapped reversibly. It was observed that SB200 ferrite/austenite interfaces are more coherent due to a finer lath thickness compared to SB300. Nevertheless, SB300 tends to trap hydrogen more strongly than SB200, as revealed by a higher T _T DA- This indicates that T _T DA decreases with interface coherency, suggesting that incoherent trapping interfaces, such as those displayed by coarse cementite, tend be irreversible.

In summary, the present invention provides for the use of a steel alloy having a microstructure comprising superbainite in a bearing component exposed to a hydrogen- containing environment to achieve a reduction in hydrogen embrittlement. It is thought that the superbainitic microstructure is able to resist hydrogen embrittlement owing to the ferrite/austenite interfaces, which may act as hydrogen traps. In other words, the reduction in hydrogen embrittlement is related to the increase in hydrogen trapping capacity of the alloy. The effect is more pronounced at higher isothermal transformation temperatures, as indicated by a comparison of SB300 and SB200. In addition, the superbainitic microstructure achieves a low hydrogen uptake compared with the conventional 10006 grade.

The foregoing description has been provided by way of explanation and illustration, and is not intended to limit the scope of the appended claims. Many variations in the presently preferred embodiments illustrated herein will be apparent to one of ordinary skill in the art, and remain within the scope of the appended claims and their equivalents.