HARD WELD OVERLAYS RESISTANT TO RE-HEAT CRACKING

Abstract

A hard weld overlay which is resistant to cracking when re-heated, and a method for designing such alloys, is disclosed. The alloys are able to resist re-heat cracking through prevention of the precipitation and/or growth of embrittling carbide, borides, or borocarbides at elevated temperatures. In one embodiment, the thermodynamics of the alloy system possess only primary carbides and secondary ferrite carbides.

Description

INCORPORATION BY REFERENCE TO ANY PRIORITY APPLICATIONS

Any and all applications for which a foreign or domestic priority claim is identified in the Application Data Sheet as filed with the present application are hereby incorporated by reference under 37 CFR 1.57.

BACKGROUND

1. Field

This invention relates to hard coatings and weld overlays used to protect surfaces from wear.

2. Description of the Related Art

The hardfacing process is a technique used to protect a surface from wear. Typical methods of hardfacing include the various methods of welding, GMAW, GTAW, PTA, laser cladding, submerged arc welding, open arc welding, thermal spray, and explosive welding. In certain applications, it is desirable for the hardfacing coating to be free of cracks. Hardbanding, the process of applying a hardfacing layer to the outer diameter of tool joints on a drill string, is an example of an application where cracks are undesirable. Cracks can allow for corrosion, create welding difficulties when re-building the hardbanding layer, and allow for the propagation of cracks from the hardfacing layer into the substrate material resulting in the failure of the drill pipe itself. Preventing cracking can be achieved in hardbanding materials by increasing the toughness of the hardfacing alloy used. However, hardness and toughness are inversely related material properties. Thus, in order to prevent cracking the hardness is sacrificed. Typical non-cracking hardfacing materials deposited via the GMAW process for the purposes of hardbanding possess hardness in the range of 50-60 HRC. Cracking hardfacing materials such as chromium carbide can exhibit hardnesses significantly above 60 HRC, in the range of 61-69 HRC.

Several modes of cracking are known to occur in hardbanding. Three types of cracking occur during welding, or slightly after (1 s-180 s) the welding as been completed. Cross checking is defined as a large crack which spans across the entire weld bead width, and can occur during the deposition of a single bead. The two other forms of cracking, dip cracking and circumferential cracking are associated with the re-heating of an existing bead. Dip cracking occurs during the welding of a single bead.

During the hardbanding process, a 1″ wide weld bead is deposited onto a rotating tool joint such that it covers the entire circumference of the joint when completed. The weld is completed when joint has made one full revolution during the weld process, such that new weld material is deposited directly on top of existing weld material. This overlap causes the existing weld material to re-heat. The re-heating of the existing weld material can cause dip cracking in the existing bead 0-2 inches away from the overlap zone.

Circumferential cracking occurs when multiple bands are welded next to each other, as is customary in the hardbanding process and other hardfacing processes. In the hardbanding process, it is customary to overlap one bead with subsequent weld passes by ⅛″ to ¼″. This slight overlap between neighboring beads re-heats the existing bead and can lead to circumferential cracking.

The thermodynamic profile of a standard hardbanding alloy, Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb_4.54Si_0.59Ti_0.39V_0.54 [alloy 1], is displayed in FIG. 1. Alloy 1 shows the solidification of the liquid (phase 9) to austenite (phase 7) before transforming to ferrite (phase 1), in addition to the solidification of several primary carbides (phase 2, 3, 4, 6, and 8) However, an important part of the thermodynamic behavior of alloy 1 is thermodynamic stability of VC (phase 4). VC is thermodynamically stable from room temperature to slightly below 1200 K, which categorizes this phase as a secondary austenite carbide. Upon initial solidification of the weld, the primary carbides and the iron matrix solidify. However, due to the rapid cooling in a weld bead, the VC phase is unlikely to form due to sluggish kinetics at these lower temperatures. The inability for VC to precipitate and grow to a stable size upon the initial deposition of the weld bead produces a super saturation of carbon in the iron-based matrix. As the weld alloy is re-heated during the continuation of the weld process, the temperature in the re-heat zone enters into the VC stable temperature range and VC can form from the ferrite or austenite. The newly formed VC introduces additional stress in the alloy and causes cracking when the weld rapidly cools again.

SUMMARY

A hard weld overlay which is resistant to cracking when re-heated, and a method for designing such alloys, is disclosed. The alloys are able to resist re-heat cracking through prevention of the precipitation and/or growth of embrittling carbide, borides, or borocarbides at elevated temperatures. Reheating of existing hard weld overlays exposes the material to elevated temperatures and occurs frequently in hard cladding operations due to overlapping and multiple pass overlays. In most hard weld overlay processes, a certain fraction of carbon and boron is trapped in the matrix due to slow cooling. Upon re-heating, the trapped carbon and boron is released to form carbides and borides. The precipitation and growth of the carbides in combination with the grain growth in the steel matrix is known to cause stress in the microstructure and lead to cracking. By controlling the thermodynamics of the boride and carbide phases, it is possible to create an alloy which is less prone to the growth of new carbides and borides during re-heating, and is thus less prone to cracking. When designing such alloys, different carbides and borides can be classified into three distinct groups: primary carbides, secondary austenite carbides, and secondary ferrite carbides.

Primary carbides are thermodynamically stable at temperatures higher than or within 50° C. of the initial solidification temperature of the matrix. Secondary austenite carbides become thermodynamically stable at temperatures above the ferrite to austenite transition temperature but more than 50° C. below the initial solidification temperature of the matrix. Finally, secondary ferrite carbides are only thermodynamically stable at temperatures below the austenite to ferrite transition.

In one embodiment, the thermodynamics of the alloy system possess only primary carbides and secondary ferrite carbides. In such as system, the primary carbides will solidify prior to the solidification of the austenite, and the remaining carbon and boron will likely be trapped in the ferrite as the alloy eventually cools to room temperature, i.e. the secondary ferrite carbides are kinetically unable to precipitate and grow to significant levels. As the weld alloy is re-heated during the continuation of the weld process, the temperature in the re-heat zone enters into the austenite region for the steel. At these temperatures, the secondary ferrite carbides are not thermodynamically stable and do not form. Grain growth of the matrix occurs, but no new significant carbide formation has occurred. As the alloy rapidly cools again, the ferrite secondary carbides are unable to form due to sluggish kinetics and the majority of the carbon and boron is once again trapped in the matrix.

In a preferred embodiment, the primary carbides are at least one of: titanium boride (TiB₂) or Niobium carbide (NbC).

Thermo-Calc is a powerful software package used to perform thermodynamic and phase diagram calculations for multi-component systems of practical importance. Calculations using Thermo-Calc are based on thermodynamic databases, which are produced by expert evaluation of experimental data using the CALPHAD method.

TCFE7 is a thermodynamic database for different kinds of steels, Fe-based alloys (stainless steels, high-speed steels, tool steels, HSLA steels, cast iron, corrosion-resistant high strength steels and more) and cemented carbides for use with the Thermo-Calc, DICTRA and TCPRISMA software packages. TCFE7 includes elements such as Ar, Al, B, C, Ca, Co, Cr, Cu, H, Mg, Mn, Mo, N, Nb, Ni, 0, P, S, Si, Ta, Ti, V, W, Zr and Fe.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1:

Phase evolution of Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb_4.54Si_0.59Ti_0.39V_0.54 [alloy 1].

FIG. 2:

Phase evolution of Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb₆Si_0.59Ti₁V_0.54 [alloy2].

FIG. 3:

Phase evolution diagram of Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb₆Si_0.59Ti₁V₂ [alloy 3].

FIG. 4: Elemental concentration in NbC phase.

FIG. 5: FCC to BCC transition temperature in selected hardbanding alloys

FIG. 6:

Phase evolution diagram of Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb_4.5Si_0.59Ti₁V_0.54.

FIGS. 7A-B: Optical microstructures at 500× of alloy 5 (7A) and alloy 6 (7B).

DETAILED DESCRIPTION

In one embodiment, the thermodynamic properties of the alloy are calculated using the CALPHAD method. A preferred embodiment uses the Thermo-Calc software to perform these calculations.

Non-Cracking Trait 1:

In one embodiment, all of the carbide, boride, and boro-carbide phases are primary carbides. Thus, they are thermodynamically stable at the relatively high temperatures as defined previously. An alloy which possesses this thermodynamic profile is more resistant to cracking than conventional hardfacing materials. As an alloy of this type is initial deposited in the form of a weld bead, the primary carbides begin to precipitate and grow during the initial solidification of the material. Typically, a large fraction of primary carbides precipitate prior to the solidification of the matrix. This solidification is advantageous for improving crack resistance, in that the existing primary carbides do not inflict high stresses on solidifying austenite or during the transformation of austenite to ferrite. The formation of primary carbides effectively reduces the total carbon in the solidifying austenite such that is less likely for the iron-based matrix to become super saturated with carbon. This aids in final structure of the metal being ferritic as opposed to austenitic, and aids in the resistance of cracking during re-heating.

In conventional hardfacing materials, the iron-based matrix is often super saturated with carbon. Upon re-heating, the carbon is allowed to diffuse throughout the microstructure and form carbides. As the matrix transforms to austenite and the grain size increases, these newly form carbides cause stresses on the microstructure of the material, which can lead to cracking in the hardfacing material.

In the described embodiment where all carbides, borides, and boro-carbide phases are primary carbides this described cracking mechanism is avoided. Upon re-heating in the described embodiment, the matrix does not form new carbides and thus stresses are avoided as the matrix transforms and grows. An example alloy [alloy 2], Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb₆Si_0.59Ti₁V_0.54, that demonstrates this phenomenon is shown in FIG. 2. This diagram shows the solidification of the liquid (phase 7) into austenite (phase 6), which ultimately transforms to ferrite (phase 2). This is the common feature of the equilibrium solidification pathway for most steels. The unique components of this alloy are the solidification of the borides, carbides, and borocarbides (phases 1, 3, 4, and 5). All of these phases can be defined as primary carbides as they form at high temperatures close to the solidification temperature of the austenite phase.

In this preferred embodiment, the primary carbides are TiB2 (phase 1), Cr2B (phase 3), NbC (phase 4), and (Fe,Cr)3B2 (phase 5).

In one embodiment the reheat temperature range is 800° C. to 1300° C.

In a preferred embodiment the reheat temperature range is 900° C. to 1200° C.

In a still preferred embodiment the reheat temperature range is 1000° C. to 1100° C.

Non-Cracking Trait 2:

In another embodiment, the mole fraction of all the carbide phases remain thermodynamically stable within the temperature range defined as the re-heat zone. In a preferred embodiment, stability is defined as a mole fraction which does not vary by more than 25%; in a still preferred embodiment stability is defined as a mole fraction which does not vary by more than 10%, in a still preferred embodiment, stability is defines as a mole fraction does not vary be more than 5%.

Carbides which are thermodynamically stable within the re-heat zone are beneficial for the purposes of creating an alloy which is resistant to re-heat cracking. In the case of a cracking prone alloy, the re-heating of the alloy can cause the precipitation and/or growth of additional carbide or the dissolution and shrinking of existing carbides. Growing or re-precipitation of carbides causes stresses in the matrix as described previously. The dissolution of carbides can also be detrimental as it increases the carbon and/or boron in the iron-based matrix. This increase in carbon in the matrix can cause other carbides to precipitate or grow causing stresses in different regions of the microstructure, or it can lead to supersaturation of carbon in the matrix which can make the material prone to re-heat cracking.

FIG. 2 depicts the thermodynamics of an alloy which possess the carbides which have a mole fraction that is thermodynamically stable within the reheat zone. As shown, there are no phase transformations or large phase mole fraction variations within the reheat zone. The primary carbide phases (1, 3, 4, and 5) are all stable from the austenite solidification temperature to temperatures below the reheat zone. When an alloy of this phase structure is re-heated, the carbides are stable and do not grow or dissolve. This prevents additional stress in the weld and cracking can be avoided.

In another embodiment, all of the secondary carbides are only thermodynamically stable below the reheat zone.

Non-Cracking Trait 3:

An alloy which possesses the thermodynamics of this embodiment is resistant to cracking in the re-heat zone. The solidification routine of such an alloy when initially deposited is similar to previously described: the Fe-based matrix and primary carbides solidify to form the microstructure. The secondary carbides are kinetically unable to form due to the rapid cooling of the process, leaving the Fe-based matrix supersaturated with carbon and/or boron. However, as the temperature of the material is increased into the reheat zone, the secondary carbide phase is not thermodynamically stable so it does not form. The material then cool rapidly down to room temperature, and the secondary carbide phase is once again unable to precipitate due to sluggish kinetics.

A preferred embodiment, Alloy Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb₆Si_0.59Ti₁V₂, is shown in FIG. 3. As shown, Phase 8, is a secondary carbide phase which is only thermodynamically stable below the reheat zone. Phase 8 is unlikely to form during the original deposition of the weld bead, and unlikely to form as the material is reheated. This embodiment allows the alloy to be supersaturated with carbon, increasing hardness, but still maintains crack resistance.

Non-Cracking Trait 4:

In another embodiment, A selection of the carbides don't contain more than 50% Fe. During reheating in the weld bead, the Fe-rich carbides can form much easily than other carbide. This phenomenon occurs because the matrix is Fe-rich and carbon has a much higher likelihood of diffusing into a region of the microstructure where Fe is free to react and precipitate new carbides. Furthermore, as the newly precipitated carbides or existing carbides are driven to grow in the alloy, the ability to utilize the large availability of Fe as opposed to lower concentration alloying elements will increase the growth rate of such carbides. Carbides which are more likely to precipitate and capable of growing rapidly in the re-heated alloy will make the alloy more susceptible to re-heat cracking.

FIG. 4 shows the variation of the mole fraction of each element in NbC, which is a common carbide in the presented hardfacing alloys. The NbC phase contains primarily Nb and C with a slight amount of V, but trace concentrations of Fe. Such a carbide will be unlikely to grow any larger during the reheating of the weld, because both Nb and V will be relatively scarce around the local region of the carbide.

In one preferred embodiment, all of the secondary carbide phases don't contain more than 50% Fe.

In a second preferred embodiment, all of the primary carbide phases don't contain more than 50% Fe.

In a still preferred embodiment, the carbide phases precipitating in the alloy consist of at least one of TiB₂, CrB₂, NbC, WC, MoB₂, and/or VC.

Non-Cracking Trait 5:

In another embodiment, the alloy is designed such that the fcc austenite/bcc ferrite transition temperature is not within the RZ. Avoiding this significant phase transformation at the RZ can minimize the stress in the microstructure and make the alloy less prone to reheat cracking. By avoiding the FCC to BCC transition upon re-heating, the alloy will be more capable of handling the stresses created by newly precipitated carbides or growth of existing carbides. FIG. 5 demonstrates how the transition temperature of the hardfacing alloy can be controlled by compositional variation.

In another embodiment, the RZ is shifted by adjusting the welding parameters used in the weld process in order to avoid the fcc austenite/bcc ferrite transition temperature in a particular alloy. The fcc austenite/bcc ferrite transition is the biggest phase transformation in the steel and can introduce significant stress causing cracking.

FIG. 5 shows the relationship between the fcc austenite/bcc ferrite transition temperature vs. carbon content. We can know what kind of microstructure (ferrite, austenite or martensite) will occur after welding by calculating the fcc austenite/bcc ferrite transition temperature. We can also adjust the fcc austenite/bcc ferrite transition temperature by changing some elements, then obtain the optimum microstructure.

Non-Cracking Trait 6:

In another embodiment, carbides do not form in the austenitic zone of the alloy during re-heating. Carbides which become stable in the austenitic zone can precipitate and/or grow upon reheating of the alloy when the matrix is austenitic. When the alloy is in the austenite phase grain growth is typical and carbides typically precipitate along the previous grain boundaries of the initially deposited ferrite matrix. Therefore, the carbides which have precipitated in the austenite are now located in the center regions of the matrix grains. As the alloy cools and transforms back to ferrite, the newly grown carbides in the center of the grains can cause stress on the microstructure and create cracks. An alloy which avoids the precipitation of carbides in the austenite zone is shown in FIG. 6. The VC, phase 3, is not thermodynamically stable in the austenite region (phase 6). Thus, any precipitation of VC do to the re-heating of the weld occurs after the alloy has transitioned from BCC to FCC upon heating and back to BCC upon cooling. Therefore, the newly formed carbide is not present during the potentially stress-inducing, and thereby crack prone, solid state transition.

In a one embodiment, the hardfacing alloy is Fe-based containing one or more of the following alloying elements B, C, Cr, Mn, Mo, Nb, Si, Ti, W, and V with additional impurities known to be present due to manufacturing procedures and possesses one of the preferred non-cracking traits described in this disclosure.

In a preferred embodiment, this hardfacing alloy is in the form of a cored welding wire.

In another preferred embodiment the hardfacing alloy composition, as defined by the composition of the feedstock material or the deposited coating, is given in weight percent by the following range: Fe_ba1C_0.5-4B_0-3Mn_0-10Al_0-5Si_0-5Ni_0-5Cr_0-30Mo_0-10V_0-10W_0-15Ti_0-10Nb_0-10

In a still preferred embodiment the hardfacing alloy composition, as defined by the composition of the feedstock material or the deposited coating, is given in weight percent by the following range: Fe_ba1C_1-2B_1-2.5Mn_1-2Al_0-0.5Si_0-1.5Ni_0-0.2Cr_0-10Mo_0-3.5V_0-2.5W_0-0.15Ti_0-2Nb_2-6

In a still preferred embodiment the hardfacing alloy composition is given in weight percent by one or a combination of the following compositions:

• Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb_4.54Si_0.59Ti_0.39V_0.54 [alloy 1]
• Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb₆Si_0.59Ti₁V_0.54 [alloy 2]
• Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb₆Si_0.59Ti₁V₂ [alloy3]
• Fe_ba1B_1.45C_0.91Cr_4.82Mn_1.01Mo_3.22Nb_4.5Si_0.59Ti_iV_0.54 [alloy 4 ]
• Fe_ba1C_1.2B₂Mn₁Si_1.1Ni_0.07Cr_8.33Mo_3.33V_0.5W_0.07Ti_1.83Nb₄ [alloy 5]
• Fe_ba1C₁B_2.5Mn₂Si_1.1Ni_0.1Cr_8.73Mo₁V_0.03W_0.03Ti_1.91Nb_4.47 [alloy 6]

EXAMPLES

Alloys 5 and 6

One of the purposes of designing alloys which possess the non-cracking traits described within this disclosure is to create a hardfacing material which exhibits very high hardness and wear resistance but is not prone to re-heat cracking. Two alloys which exhibit both high hardness and resistance to re-heat cracking are alloys 5 and 6. Alloys 5 and 6 where produced in the form of welding wires and welded onto a standard 6⅝″ O.D. tool joint in a manner customary to the hardband process used in the oil and gas industry. The feedstock wires were also melted into small ingots in an arc-melter, for the purposes of measuring un-diluted hardness and examining microstructure. The results of the hardness measurements for both ingot form and weld bead form are shown in Table 1. Both alloys exhibit high hardness 60 HRC or above, a region which is not typical for crack resistant hardfacing alloys.

TABLE 1
Hardness values of selected disclosed alloys
Alloy	Form	Hardness
5	Ingot	63-67
5	Weld Bead	61-63
6	Ingot	59-60

The microstructures of alloy 5 and 6 are shown in FIGS. 7A-B. Both alloys show a high frequency of carbides within the microstructure which provides good hardness and wear resistance, but is typically an indicator for the alloy being prone to cracking. However, both alloys were deposited via a process typically used in hardbanding as three consecutive bands and were free of any cracks. The hardbanding process used reheats existing bead deposits, and is known to generate both dip cracks and circumferential cracks in crack prone alloys of lesser hardness.

Claims

1. (canceled)

2. A hardfacing weld deposit comprising: a hardness of at least 60 HRC; anda microstructure comprising: an iron-based matrix; andcarbides and/or borides;wherein the carbides and/or borides comprise only carbides and/or borides which precipitate prior to solidification of the iron-based matrix.

3. The deposit of claim 2, wherein the carbides and/or borides are selected from the group consisting of titanium boride, niobium carbide, chromium boride, iron-chromium boride, and combinations thereof

4. The deposit claim 2, wherein the deposit does not form additional carbides or borides when re-heated to a range of 800° C. to 1300° C. for 1 s to 180 s.

5. The deposit of claim 2, wherein the deposit does not form additional carbides or borides when re-heated to a range of 900° C. to 1200° C. for 1 s to 180 s.

6. The deposit of claim 2, wherein the deposit does not form additional carbides or borides when re-heated to a range of 1000° C. to 1100° C. for 1 s to 180 s.

7. The deposit of claim 2, wherein the deposit comprises at least one of:

8. A hardfacing weld deposit comprising: a hardness of at least 60 HRC; anda stable carbide and/or boride structure;wherein a mole fraction of the stable carbide and/or boride structure does not change by more than 25% when reheated.

9. The deposit of claim 8, wherein the stable carbide and/or boride structure in the deposit does not change when re-heated to a range of 800° C. to 1300° C. for 1 s to 180 s.

10. The deposit of claim 8, wherein the mole fraction of the stable carbide and/or boride structure does not change by more than 10% when reheated.

11. The deposit of claim 8, wherein the mole fraction of the stable carbide and/or boride structure does not change by more than 5% when reheated.

12. The deposit of claim 8, wherein the deposit further comprises an iron-based matrix, and the deposit possesses a carbide and/or boride thermodynamic stability such that a mole fraction of the carbides and/or borides does not change by more than 25% over a temperature range between room temperature and a solidification temperature of the iron-based matrix.

13. The deposit of claim 8, wherein the deposit further comprises an iron-based matrix, and the deposit possesses a carbide and/or boride thermodynamic stability such that any carbides and/or borides do not form at temperatures above the solidification temperature of the iron-based matrix, and are only stable at temperatures below a re-heat temperature range.

14. The deposit of claim 13, wherein the re-heat temperature range is about 800° C. to 1300° C.

15. The deposit of claim 13, wherein the re-heat temperature range is about 900° C. to 1200° C.

16. The deposit of claim 13, wherein the re-heat temperature range is about 1000° C. to 1100° C.

17. The deposit of claim 8, wherein the deposit comprises at least one of:

18. A hardfacing weld deposit comprising: a hardness of at least 60 HRC; andcarbides and/or borides;wherein the carbides and/or borides comprise an iron concentration of 50 wt. % or less.

19. The deposit of claim 18, wherein the carbides and/or borides are selected from the group consisting of niobium carbide, titanium boride, chromium boride, tungsten carbide, molybdenum boride, and vanadium carbide, and combinations thereof.

20. A hardfacing weld deposit comprising: a hardness of at least 60 HRC; andan austenite to ferrite transition temperature which is outside a re-heat temperature range.

21. The deposit of claim 20, wherein the re-heat temperature range is about 800° C. to 1300° C.

22. The deposit of claim 20, wherein the re-heat temperature range is about 900° C. to 1200° C.

23. The deposit of claim 20, wherein the re-heat temperature range is about 1000° C. to 1100° C.

24. The deposit of claim 20, wherein the deposit comprises at least one of: