metals

Review Strengthening Mechanisms in Nickel-Copper Alloys: A Review

Olexandra Marenych and Andrii Kostryzhev * School of Mechanical, Materials, Mechatronic and Biomedical Engineering, University of Wollongong, Wollongong, NSW 2500, Australia; [email protected] * Correspondence: [email protected]; Tel.: +61-02-4221-3034



Received: 2 September 2020; Accepted: 30 September 2020; Published: 12 October 2020


Abstract: Nickel-Copper (Ni-Cu) alloys exhibit simultaneously high strength and toughness, excellent corrosion resistance, and may show good wear resistance. Therefore, they are widely used in the chemical, oil, and marine industries for manufacturing of various components of equipment, such as: drill collars, pumps, valves, impellers, fixtures, pipes, and, particularly, propeller shafts of marine vessels. Processing technology includes bar forging, plate and tube rolling, wire drawing followed by heat treatment (for certain alloy compositions). Growing demand for properties improvement at a reduced cost initiate developments of new alloy chemistries and processing technologies, which require a revision of the microstructure-properties relationship. This work is dedicate to analysis of publicly available data for the microstructure, mechanical properties and strengthening mechanisms in Ni-Cu alloys. The effects of composition (Ti, Al, Mn, Cr, Mo, Co contents) and heat treatment on grain refinement, solid solution, precipitation strengthening, and work hardening are discussed.

Keywords: Ni-Cu alloys; microstructure characterisation; mechanical properties; strengthening mechanisms

1. Introduction The Ni-Cu system forms the basis for the Monel alloy family (Table1). Monel was discovered by Robert Crooks Stanley who was employed by the International Nickel Company (INCO) in 1901. The new alloy was named in honour of the company president, Ambrose Monell. The name is now a trademark of Special Metals Corporation [1].

Table 1. Chemical compositions of Monel alloys (wt.%).

Monel Alloys Ni Cu C Mn Fe Co S Si Al Ti Max - 34.0 0.3 2.0 2.5 - 0.024 0.5 - - Monel 400 Min 63.0 28.0 ------Max 45.0 Balance 0.1 2.25 0.75 0.25 0.015 0.25 - - Monel 401 Min 40.0 ------Max 57.0 Balance 0.15 0.1 0.5 - 0.024 0.1 0.05 - Monel 404 Min 52.0 ------Max - 34.0 0.3 2.0 2.5 - 0.060 0.5 - - Monel R405 Min 63.0 28.0 - - - - 0.025 - - - Max - 33.0 0.18 1.5 2.0 0.25 0.006 0.5 3.15 0.85 Monel K500 Min 63.0 27.0 ------2.3 0.35

Ni and Cu exhibit very similar atomic characteristics. They both have face centred cubic (fcc) crystal structure type, less than three percent difference in atomic radii, and exhibit similar electronegativity

Metals 2020, 10, 1358; doi:10.3390/met10101358 www.mdpi.com/journal/metals Metals 2020, 10, 1358 2 of 18 and valence state. The Ni-Cu system has complete solid solubility, which allows production of single phase alloys over the entire composition range [2]. Although the Ni-Cu system exhibits complete solid solubility [3], the large differences in melting points between Ni (1455 ◦C) and Cu (1085 ◦C) can result in Cu segregation. Following equilibrium solidification at slow cooling rates, dendrites become enriched in Ni and interdendritic regions get enriched in Cu [4–6]. However, with an increase in cooling rate during solidification the compositional gradient decreases and the microstructure morphology changes from dendritic to cellular [7]. Higher undercoolings during solidification also lead to finer and more equiaxed grain sizes after annealing [8]. Monel alloys can be easily fabricated by hot and cold metal forming processes and machining. Recrystallisation studies determined the optimum hot deformation temperatures to be 950–1150 ◦C[9,10], which are quite similar to other Ni-base alloys and steels. However, heat treatment schedule requires rigorous development: usually a two-step age-hardening heat treatment in the temperature range of 650–480 ◦C is used for Ni-Cu alloys [11,12]. The higher temperature stage helps to quickly nucleate precipitates of alloying elements, and the lower temperature stage provides a superior distribution of higher number density of smaller-sized particles. Ni-Cu alloys are usually quite weldable to each other and to other Ni alloys and stainless steels [13,14]. Lower heat inputs produce finer grain microstructures with random texture and higher strength and ductility [15]. Monel alloys are expensive, with their cost reaching up to 3 times that of Ni and 7 times that of Cu [16–18]. Hence their use is limited to those applications where they cannot be replaced with a cheaper alternative. Major additions of copper (28–40 wt.%) improve corrosion resistance of Ni in many agents, in particular nonoxidizing acids, nonaerated sulphuric and hydrofluoric acids [19–21]. This determines areas of application of Ni-Cu alloys. They are widely used for manufacturing various components of equipment in chemical, oil and marine industries (such as drill collars, pumps, valves, fixtures, piping, fasteners, screws, propeller shafts, steam generators, turbines [22–27]), for protective coating [28,29], for manufacturing electrical and electronic equipment (resistors, bimetal contacts, capsules for transistors and ceramic-to-metal sealing [30,31]), and in fuel cells [32,33].

2. Role of Alloying Elements Ni can dissolve high concentrations of alloying elements compared to other metals [34]. Alloying elements increase mechanical strength via solid solution [35–40] or precipitation strengthening [41–45]. However, the solid solution strengthening effect of Cu is much lower compared to other elements [19,46]. For example, in recently developed nanocrystalline Ni-Cu alloys both strength and ductility increased with Cu content by only 40% for Cu concentrations increasing by 5 times (from 6 to 32 wt.%) [47]. This was associated with grain refinement and solid solution strengthening. Work hardening also increased with Cu content, as the stacking fault energy decreased, leading to increased dislocation densities and twinning. Due to the substantial increase in strength-to-weight ratio, achieved by precipitation hardening, the weight of turbine engines was significantly reduced, which enhanced development of the aerospace industry [22]. Although Ti and Al provide some solid solution strengthening effect, they typically improve strength by precipitation of γ0-Ni3(Ti,Al) particles during heat treatment [19,48]. Aluminium (2.3–3.15%) provides excellent corrosion resistance resulting from the formation of a protective Al2O3 surface oxide layer. Titanium (0.35–0.85%) forms TiC carbides [49]. Together Al and Ti are often used in minor amounts to increase corrosion resistance via deoxidation [20]. They both combine with oxygen to form oxides, thus controlling porosity in welds [2]. Carbon (<0.25 wt.%) is required for carbide formation, which not only increase strength at room temperature but also creep resistance [50]. Manganese (<1.5%) improves corrosion resistance and weldability, promotes formation of M23C6 type carbides [3,51]. Iron (<2.0%) provides solid solution strengthening at reduced costs, but may be detrimental for corrosion resistance [20]. Fe also increases the solubility of C in Ni; this improves resistance to high-temperature carburizing environments [19]. Cobalt (<0.25%) increases high-temperature strength via solid solution strengthening, and resistance to carburization and sulphidation [20]. Additions of Co also raise solvus temperature of γ0 phase [19]. Sulphur Metals 2020, 10, 1358 3 of 18 Metals 2020, 10, x FOR PEER REVIEW 3 of 19

(raise<0.006%) solvus enhances temperature machinability of γ′ phase [52 ]. [19 Silicon]. Sulphur (<0.5%) (<0.006%) is typically enhances present machinability only in minor [52 amounts]. Silicon as(<0.5%) a residual is typically element present from deoxidationonly in minor or amounts as an intentional as a residual addition element to promote from deoxidation high-temperature or as an oxidationintentional resistance addition [to19 ].promote high-temperature oxidation resistance [19].

3.3. Microstructure Ni-CuNi-Cu alloysalloys havehave fccfcc matrixmatrix ((γγ)) withwith aa latticelattice parameterparameter ofof 0.35340.3534 nmnm [[5353].]. DueDue toto fullfull mutualmutual solubilitysolubility ofof NiNi andand Cu,Cu, thethe secondsecond phasephase doesdoes notnot formform (Figure(Figure1 ),1), although although particles particles of of various various chemistrieschemistries can precipitate. precipitate. Therefore, Therefore, in inthis this alloy alloy strengthening strengthening occurs occurs via four via mechanisms: four mechanisms: grain grainrefinement, refinement, solid solutionsolid solution,, precipitation precipitation and dislocation and dislocation strengthening strengthening (work (work hardening hardening).). Grain Grainrefinement refinement is achieved is achieved as a asresult a result of recrystallization of recrystallization during during hot hot deformation deformation in in the the 87 870–11500–1150 ◦°CC temperaturetemperature rangerange [[5454].]. AA requiredrequired levellevel ofof precipitationprecipitation strengtheningstrengthening isis usuallyusually obtainedobtained afterafter ageage hardeninghardening heatheat treatment.treatment. DislocationDislocation strengtheningstrengthening followsfollows coldcold rollingrolling oror drawingdrawing [[5555––5757].].

FigureFigure 1.1. TypicalTypicalmicrostructure microstructure of of Monel Monel K500 K500 rod rod in in (a )(a hot) ho rolledt rolled and and (b )(b annealed) annealed at 1100at 1100◦C forC for 15 min15 min conditions conditions (the (the authors’ authors’ unpublished unpublished data). data).

AfterAfter an an appropriate appropriate heat heat treatment, treatment, the following the following intermetallic intermetallic particles particles can precipitate: can precipitate Ni3(AlX),: NiNi33Ti,(Al andX), Ni NiFe3Ti,3 and(AlFe) Ni [Fe493,(Al58].Fe) Ni [349(Al,X),58]. Ni particles,3(Al,X) particles, where X canwhere be Cu,X can Mn, be Ti, Cu, or Mn, Si, are Ti, known or Si, are as gammaknown as prime gamma precipitates prime precipitates (γ’). Gamma (γ’). prime Gamma (γ’) prime particles (γ’) have particles an ordered have an fcc ordered crystal fcc structure crystal (Figurestructure2a) (Figure with close 2a) with crystallographic close crystallographic matching matching to the Ni fccto the matrix Ni fcc ( < matrix1% mismatch) (<1% mismatch [59]. The) [59 Al]. atomsThe Al occupy atoms occupy the cube the corners cube corners while thewhile Ni the ones Ni resideones reside on the on face the centres.face centres. In Ni-Cr-Fe In Ni-Cr alloys-Fe alloys Co andCo and Fe canFe can substitute substitute Ni Ni [60 [],60 while], while Cu, Cu, Mn, Mn and, and Si Si tend tend to to substitute substitute Al. Al. Gamma Gamma primeprime particlesparticles homogeneouslyhomogeneously nucleatenucleate throughthrough thethe volume,volume, andand maintainmaintain sphericalspherical shapeshape andand coherencycoherency withwith thethe NiNi matrixmatrix eveneven afterafter manymany hourshours ofof agingaging (Figure(Figure3 3).). The The precipitates precipitates do do not not show show any any tendency tendency for for alignmentalignment oror nucleationnucleation atat dislocations.dislocations. CloseClose crystallographiccrystallographic matchingmatching betweenbetweenthe theγ γmatrix matrix andandγ γ′0 precipitatesprecipitates leads leads to to a very a very low surfacelow surface energy energy that, consequently, that, consequently results, in results very low in particlevery low coarsening particle ratescoarsening and increased rates and number increased density number of fine densityparticles of [61 fine]. The particles coarsening [61] process. The coarsening is controlled process by bulk is dicontrolledffusion of by the bulkγ’-forming diffusion elements. of the Theγ’-forming activation elements energy. ofThe the activation coarsening energy process of is the close coarsening to that of diprocessffusion is of close Al in to Ni. that The of averagediffusionγ ’of particle Al in Ni. radius The remainsaverage proportionalγ’ particle radius to the remains cube root proportional of ageing time.to the The cube volume root of fraction ageing time. of precipitates The volume can fraction reach 6–7% of precipitates after aging can at 700reach◦C[ 6–497%]. after aging at 700 °C [49The]. Ni-Al and Ni-Ti phase diagrams (Figure4) can help to analyse precipitation in Ni-base alloys [62]. Most Ni-base alloys contain below 10 wt.% of (Al + Ti). However, even minor additions of Al or Ti result in precipitation of Ni3Al or Ni3Ti particles. The γ’ solvus temperature is ranging between 700 and 750 ◦C for the solution treated and quenched alloys [49]. Therefore, the solution annealing temperature is usually above 950 ◦C. The solubility temperature of γ’ particles in Monel is relatively lower compared with those in high temperature precipitation strengthened Ni-base superalloys. This is associated with much lower Al and Ti contents in Monel K500 compared to Al, Ti, Nb, and Ta concentrations in other Ni-base alloys [63]. Decomposition of the supersaturated austenite may not occur during quenching, meaning that even very fine nano-scale particles do not precipitate [49]. This is contrarily to other Ni-Al and Ni-Ti alloys, where γ’ forms readily on Metals 2020, 10, 1358 4 of 18

quenching [64,65]. The substitution of some neighbouring elements in the periodic table (Cr, Fe, and Co) for Ni in alloy composition is known to decrease γ’ solubility and increase the γ’-particle volume fraction even for the same Al + Ti contents [66]. Precipitation of γ’ in Monel K500 at 750 ◦C, while the solubility of Al in Ni at this temperature is quite high and is reaching 5 wt.% in binary Ni-A1 system, suggests that Cu also decreases the solubility of γ’[67]. The γ0 precipitation provides significant strengthening effect with temperatures increasing up to ~800 ◦C[68]. This is believed to occur due to ordering in the particle lattice and the relatively low mobility of super lattice dislocations Metals 2020, 10, x FOR PEER REVIEW 4 of 19 thatMetals occurs 2020, 10 with, x FOR increasing PEER REVIEW temperature. 4 of 19

Figure 2. (a) Ordered face centred cubic (fcc) crystal structure of a γ′-Niγ 3(Ti,Al) particle and (b) TiC Figure 2. (a(a) )Ordered Ordered face face centred centred cubic cubic (fcc (fcc)) crystal crystal structure structure of ofa γ′ a-Ni0-Ni3(Ti,Al)3(Ti,Al) particle particle and and(b) TiC (b) particle. TiCparticle. particle.

FigureFigure 3. TEMTEM dark field imagesimages illustratingillustrating thethe homogeneoushomogeneous distribution distribution and and the the spherical spherical shape shape of Figure 3. TEM dark field images illustrating the homogeneous distribution and the spherical shape ofthe theγ γ′precipitates precipitates retained retained even even after after prolonged prolonged aging aging at 700 at 700C for °C ( afor) 64 (a h,) 64 (b )h, 128 (b h,) 128 and h (,c and) 1000 (c) h 1000 [49]. of the0 γ′ precipitates retained even after prolonged aging at◦ 700 °C for (a) 64 h, (b) 128 h, and (c) 1000 Metalsh 2020 [49],. 10 , x FOR PEER REVIEW 5 of 19 h [49]. The Ni-Al and Ni-Ti phase diagrams (Figure 4) can help to analyse precipitation in Ni-base The Ni-Al and Ni-Ti phase diagrams (Figure 4) can help to analyse precipitation in Ni-base alloys [62]. Most Ni-base alloys contain below 10 wt.% of (Al + Ti). However, even minor additions alloys [62]. Most Ni-base alloys contain below 10 wt.% of (Al + Ti). However, even minor additions of Al or Ti result in precipitation of Ni3Al or Ni3Ti particles. The γ’ solvus temperature is ranging of Al or Ti result in precipitation of Ni3Al or Ni3Ti particles. The γ’ solvus temperature is ranging between 700 and 750 °C for the solution treated and quenched alloys [49]. Therefore, the solution between 700 and 750 °C for the solution treated and quenched alloys [49]. Therefore, the solution annealing temperature is usually above 950 °C. The solubility temperature of γ’ particles in Monel is annealing temperature is usually above 950 °C. The solubility temperature of γ’ particles in Monel is relatively lower compared with those in high temperature precipitation strengthened Ni-base relatively lower compared with those in high temperature precipitation strengthened Ni-base superalloys. This is associated with much lower Al and Ti contents in Monel K500 compared to Al, superalloys. This is associated with much lower Al and Ti contents in Monel K500 compared to Al, Ti, Nb, and Ta concentrations in other Ni-base alloys [63]. Decomposition of the supersaturated Ti, Nb, and Ta concentrations in other Ni-base alloys [63]. Decomposition of the supersaturated austenite may not occur during quenching, meaning that even very fine nano-scale particles do not austenite may not occur during quenching, meaning that even very fine nano-scale particles do not precipitate [49]. This is contrarily to other Ni-Al and Ni-Ti alloys, where γ’ forms readily on precipitate [49]. This is contrarily to other Ni-Al and Ni-Ti alloys, where γ’ forms readily on quenching [64,65]. The substitution of some neighbouring elements in the periodic table (Cr, Fe, and quenching [64,65]. The substitution of some neighbouring elements in the periodic table (Cr, Fe, and Co) for Ni in alloy composition is known to decrease γ’ solubility and increase the γ’-particle volume Co) for Ni in alloy composition is known to decrease γ’ solubility and increase the γ’-particle volume fraction even for the same Al + Ti contents [66]. Precipitation of γ’ in Monel K500 at 750 °C, while the fraction even for the sameFigure Al 4.+ TiBinary contents ( a) Ni Ni-Al [-6Al6]. and Precipitation ( b) Ni Ni-Ti-Ti phase of γ diagramsdiagram’ in Monels [[626 2K500].]. at 750 °C, while the solubility of Al in Ni at this temperature is quite high and is reaching 5 wt.% in binary Ni-A1 system, solubility of Al in Ni at this temperature is quite high and is reaching 5 wt.% in binary Ni-A1 system, suggestsThe Nithat3Ti Cu phase, also designateddecreases the as eta sol (ubilityη), exhibits of γ’ stoichiometric [67]. The γ′ precipitation composition providesof 75 at.% significant Ni-25 at.% suggestsThe Nithat3Ti Cuphase, also designated decreases asthe eta sol (η),ubility exhibits of γ stoic’ [67hiometric]. The γ′ precipitationcomposition of provides 75 at.% N significanti-25 at.% strengtheningTi and a stable effect hexagonal with temperatures closed packed increasing crystal structure. up to ~800 Although °C [68]. This the η is-Ni believed3Ti particles to occur may due have to Tistrengthening and a stable effect hexagonal with temperatures closed packed increasing crystal structure. up to ~800 Although °C [68]. theThis η is-Ni believed3Ti particles to occur may due have to orderingordered metastablein the particle fcc structurelattice and at the lower relatively temperatures. low mobility Ti additions of super promote lattice formationdislocations of that the η occurs-Ni Ti orderedordering metastable in the particle fcc latticestructure and at the lower relatively temperatures. low mobility Ti additions of super latticepromote dislocations formation that of occursthe 3η- withphase, increasing while Al temperature. additions promote formation of the γ0-Ni3Al phase. The hexagonal η-Ni3Ti usually Niwith3Ti increasing phase, while temperature. Al additions promote formation of the γ′-Ni3Al phase. The hexagonal η-Ni3Ti usually appears as coarse disk-shaped particles of low number density that does not provide reasonable strengthening. Thus, fine γ′-Ni3Al particles of high number density are preferred [68]. Appreciable C levels, 0.18 wt.% in Monel K500, result in precipitation of carbides. Carbides composition, size, and number density depend on the alloy composition and processing history, and may vary in service [2]. The Ti-rich MC carbide is the main carbide phase in Monel K500, (Figure 2b). Coarse TiC particles randomly precipitate within the grains or along the grain boundaries (Figure 5) [69]. They possess the fcc crystal structure and usually form at the end of solidification by the eutectic- type reaction [70]. The morphology and distribution of TiC do not change with heat treatment below 980 °C. However, these particles may slightly grow during heat treatment [69]. TiC do not easily dissolve during re-heating for metal forming operations and, thus, may appear as stringers along the rolling direction. The MC carbides can often be replaced by M23C6 carbides during thermomechanical processing or high temperature service [48]. Due to absence of Cr, the M23C6 carbides in Monel K500 are generally Mn rich and precipitate in the temperature range of 760–980 °C . These carbides preferentially form on grain boundaries and, thus, can improve creep resistance via restricting grain boundary migration (Figure 5c). Presence of carbides retards recrystallization, refines grain size, and improves the high temperature strength and ductility. However, the carbides type and morphology should be carefully controlled to avoid fracture initiation at the particle-matrix interface during a component manufacturing or in service [71]. The MCN carbonitrides are similar to MC carbides in morphology and crystallography, except that substantial levels of C are substituted by N. The eutectic-type reactions leading to formation of carbides and carbonitrides are facilitated by the strong tendency of C, N, and Ti to segregate to the liquid during solidification. As a result, carbonitrides may precipitate in the interdendritic regions and grain boundaries [71,72]. Recent studies of boron addition to Monel 400 showed a possibility to increase hardness by 4 times (from 200 HV to 800 HV) and wear resistance by 60% following precipitation of Ni borides [73,74].

Metals 2020, 10, 1358 5 of 18 appears as coarse disk-shaped particles of low number density that does not provide reasonable strengthening. Thus, fine γ0-Ni3Al particles of high number density are preferred [68]. Appreciable C levels, 0.18 wt.% in Monel K500, result in precipitation of carbides. Carbides composition, size, and number density depend on the alloy composition and processing history, and may vary in service [2]. The Ti-rich MC carbide is the main carbide phase in Monel K500, (Figure2b). Coarse TiC particles randomly precipitate within the grains or along the grain boundaries (Figure5)[ 69]. They possess the fcc crystal structure and usually form at the end of solidification by the eutectic-type reaction [70]. The morphology and distribution of TiC do not change with heat treatment below 980 ◦C. However, these particles may slightly grow during heat treatment [69]. TiC do not easily dissolve during re-heating for metal forming operations and, thus, may appear as stringers along the rolling direction. The MC carbides can often be replaced by M23C6 carbides during thermomechanical processing or high temperature service [48]. Due to absence of Cr, the M23C6 carbides in Monel K500 are generally Mn rich and precipitate in the temperature range of 760–980 ◦C. These carbides preferentially form on grain boundaries and, thus, can improve creep resistance via restricting grain boundary migration (Figure5c). Presence of carbides retards recrystallization, refines grain size, and improves the high temperature strength and ductility. However, the carbides type and morphology should be carefully controlled to avoid fracture initiation at the particle-matrix interface during a component manufacturing or in service [71]. The MCN carbonitrides are similar to MC carbides in morphology and crystallography, except that substantial levels of C are substituted by N. The eutectic-type reactions leading to formation of carbides and carbonitrides are facilitated by the strong tendency of C, N, and Ti to segregate to the liquid during solidification. As a result, carbonitrides may precipitate in the interdendritic regions and grain boundaries [71,72]. Recent studies of boron addition to Monel 400 showed a possibility to increase hardness by 4 times (from 200 HV to 800 HV) and wear resistance by 60% following precipitation of Ni borides [73,74]. Metals 2020, 10, x FOR PEER REVIEW 6 of 19

FigureFigure 5.5. PParticlearticle precipitation: ( (aa)) TiC TiC within the the grains, ( (bb)) TiC at the grain boundaries,boundaries, and (c)

MM2323CC66-type-type carbides carbides at atthe the grain grain boundaries boundaries [69]. [ 69].

AA moderatemoderate densitydensity ofof dislocationsdislocations havehave beenbeen observedobserved inin MonelMonel K500K500 thinthin (0.2(0.2 mm)mm) sheetssheets thatthat werewere solutionsolution annealed,annealed, aged,aged, andand waterwater quenchedquenched [[4949].]. ItIt waswas proposedproposed thatthat thethe dislocationsdislocations werewere formedformed followingfollowing micro-deformationsmicro-deformations underunder thethe quenchingquenching stresses.stresses. These dislocationsdislocations ofof moderatemoderate densitydensity diddid notnot exhibitexhibit anyany distributiondistribution patternpattern inin thethe formform ofof planarplanar arraysarrays oror pairing.pairing. This suggestssuggests aa quitequite highhigh levellevel ofof stackingstacking faultfault energyenergy ofof thethe alloyalloy andand absenceabsence ofof short-rangeshort-range orderingordering inin thethe solution-annealedsolution-annealed and quenchedquenched material.material. The arrangement of dislocations in thethe solution-treatedsolution-treated andand deformed deformed (3%(3% strain)strain) specimens specimens was found to be be nonplanar nonplanar in several several regions regions of of the the grains grains (Figure(Figure6 a)6a [) 49[49].]. Although Although the the stacking stacking fault fault energy energy of of Ni Ni alloys alloys is highis high [75 ],[75 alloying], alloying may may decrease decrease this value,this value, as has as beenhas been observed, observed for, example,for example in many, in many Ni-Cr-base Ni-Cr-base alloys alloys [76 ].[76 However,]. However, Cu Cu additions additions to Nito Ni do do not not reduce reduc thee the stacking stacking fault fault energy energy of of the the alloy alloy [77 [77,78,78].]. The The concentrations concentrations ofof otherother alloyingalloying elementselements inin MonelMonel K500 K500 are are too too low low to to substantially substantially reduce reduc thee the stacking stacking fault fault energy energy of the of matrixthe matrix [49]. Therefore,[49]. Therefore, deformation deformation twins are twins usually are notusually observed. not observed. However, However, annealing annealingtwins do form twins (Figure do form6b). Annealing(Figure 6b). twins Annealing in Monel twins K500 in are Monel prevalent K500 inare solution prevalent annealed, in solution water annealed quenched,, water air-cooled, quenched, or aged air- materialcooled, or [49 aged,72]. material The solution [49,72 annealed]. The solution alloy annealed was found alloy to containwas found a moderate to contain density a moderate of annealing density twins,of annealing which twins, usually which appeared usually as parallelappeared bands as parallel bounded bands by coherent bounded (111) by coherent planes. (111) planes.

Figure 6. (a) Non-planar dislocation arrangement in a slightly deformed (3% strain) sample [49] and (b) annealing twins in the fully annealed Monel K500 [69].

4. Strengthening Mechanisms As Monel K500 is a single phase alloy, four strengthening mechanisms operate in it: grain refinement, solid solution, precipitation, and dislocation strengthening (work hardening). Depending on the microalloying element content (in particular, Ti, Al, Co, Fe) and processing schedule (age hardening heat treatment), either the solid solution strengthening or precipitation strengthening dominates.

4.1. Grain Refinement Figure 7a shows an example of grain size strengthening in Monel 400. With an increase in grain size from 9.5 to 202 µm the stress-strain curve moves downwards along the stress axis and the yield stress value decreases from about 290 to 130 MPa. However, this variation in grain size did not affect the work-hardening behaviour, i.e., the stress-strain curves for all the grain sizes are essentially Metals 2020, 10, x FOR PEER REVIEW 6 of 19

Figure 5. Particle precipitation: (a) TiC within the grains, (b) TiC at the grain boundaries, and (c)

M23C6-type carbides at the grain boundaries [69].

A moderate density of dislocations have been observed in Monel K500 thin (0.2 mm) sheets that were solution annealed, aged, and water quenched [49]. It was proposed that the dislocations were formed following micro-deformations under the quenching stresses. These dislocations of moderate density did not exhibit any distribution pattern in the form of planar arrays or pairing. This suggests a quite high level of stacking fault energy of the alloy and absence of short-range ordering in the solution-annealed and quenched material. The arrangement of dislocations in the solution-treated and deformed (3% strain) specimens was found to be nonplanar in several regions of the grains (Figure 6a) [49]. Although the stacking fault energy of Ni alloys is high [75], alloying may decrease this value, as has been observed, for example, in many Ni-Cr-base alloys [76]. However, Cu additions to Ni do not reduce the stacking fault energy of the alloy [77,78]. The concentrations of other alloying elements in Monel K500 are too low to substantially reduce the stacking fault energy of the matrix [49]. Therefore, deformation twins are usually not observed. However, annealing twins do form (Figure 6b). Annealing twins in Monel K500 are prevalent in solution annealed, water quenched, air- Metalscooled,2020 or, 10aged, 1358 material [49,72]. The solution annealed alloy was found to contain a moderate density6 of 18 of annealing twins, which usually appeared as parallel bands bounded by coherent (111) planes.

Figure 6. (a) Non-planarNon-planar dislocationdislocation arrangement in a slightly deformed (3% strain) sample [49] andand ((b)) annealingannealing twinstwins inin thethe fullyfully annealedannealed MonelMonel K500K500 [[6969].].

4. Strengthening M Mechanismsechanisms As Monel Monel K500 K500 is is a a single single phase phase alloy, alloy, four four strengthening strengthening mechanisms mechanisms operate operate in it: grain in it: grainrefinement, refinement, solid solution, solid solution, precipitation precipitation,, and dislocation and dislocation strengthening strengthening (work hardening) (work. hardening).Depending Dependingon the microalloying on the microalloying element content element (in particular, content (in Ti, particular, Al, Co, Fe) Ti, and Al, processing Co, Fe) and schedule processing (age schedulehardening (age heat hardening treatment) heat, either treatment), the solid either solution the solid strengthening solution strengthening or precipitation or strengthening precipitation strengtheningdominates. dominates.

4.1. Grain Refinement 4.1. Grain Refinement FigureFigure7 7aashows shows anan exampleexample ofof graingrain sizesizestrengthening strengtheningin inMonel Monel400. 400. WithWith anan increaseincreasein ingrain grain size from 9.5 to 202 µm the stress-strain curve moves downwards along the stress axis and the yield Metalssize from2020, 10 9.5, x toFOR 202 PEER µm REVIEW the stress -strain curve moves downwards along the stress axis and the7 yieldof 19 stress value decreases fromfrom about 290 to 130 MPa. However, this variation in grain size did not a ffectffect parallel.the work-hardening work Independence-hardening behaviour, behaviour, of work i.e.,hardening i.e. the, the stress-strain stressrate on-strain grain curves curves was for observed all for the all grain thein other sizes grain arefcc sizes essentiallymetals, are especially essentially parallel. atIndependence low initial dislocation of work densities, hardening when rate ona possible grain was dislocation observed accumulation in other fcc at metals, grain boundaries especially at does low notinitial restrict dislocation the dislocation densities, motion when in a the possible grain interior dislocation [79]. accumulation With strain rate at grainincreasing boundaries from the does quasi not- staticrestrict to thedynamic dislocation, the stress motion-strain in the curves grain interiormove up [79 along]. With the strain stress rate axis increasing (Figure from7b); a the similar quasi-static trend wasto dynamic, observed the in other stress-strain fcc materials curves [80 move]. However, up along the the effect stress of axisstrain (Figure rate on7b); the a stress similar-strain trend curve was dependenceobserved in on other grain fcc size materials is weak: [ 80the]. wo However,rk hardening the eraffectte at of a high strain strain rate rates on the remains stress-strain invariant curve of graindependence size [81 on]. grain size is weak: the work hardening rate at a high strain rates remains invariant of grain size [81].

Figure 7. The stress-strain curves of Monel 400 for various grain sizes at (a) low and (b) high strain rates [81]. Figure 7. The stress-strain curves of Monel 400 for various grain sizes at (a) low and (b) high strain rates [81].

4.2. Solid Solution Strengthening The solid solution strengthening effect is known to increase with the atomic size difference between the primary element and a solute. Addition of the substitutional alloying elements results in expansion of the fcc lattice leading to strengthening [46,82]. Table 2 summarizes approximate atomic diameter variations between Ni and some other elements and solubility data for these elements in Ni at 1000 °C. As can be seen, Ta, W, Mo, Ti, V, Cr, and Mn provide the best combination of atomic radii mismatch and appreciable solubility needed for solid solution strengthening. Therefore, Mo, Ti, Cr, and Mn are frequently used in many single phase Ni-base alloys for solid solution strengthening. Fe, Co and Cu also can be used for solid solution strengthening due to their high solubility in Ni [83]. The solid solution strengthening from additions of these elements will be effective only if the element concentrations do not exceed their solubility limits in Ni-based austenite, otherwise particle precipitation will take place. Although Ti, Al, and Nb can be effective solid solution strengtheners, they typically improve strength via precipitation of γ′-Ni3(Ti,Al) or Ni3Nb intermetallic particles, or carbides [2]. In addition to the atomic size misfit, electron configuration [84], diffusion coefficients [39], and γ-γ′ lattice misfit (dependent on element partitioning between γ- matrix and γ′-particles) [85] were shown to influence the solid solution strengthening at high temperatures. Thus, slow diffusing elements Re, Os, Ir, Tc, Ru, and W provide the best high temperature strength and creep resistance.

Metals 2020, 10, 1358 7 of 18

4.2. Solid Solution Strengthening The solid solution strengthening effect is known to increase with the atomic size difference between the primary element and a solute. Addition of the substitutional alloying elements results in expansion of the fcc lattice leading to strengthening [46,82]. Table2 summarizes approximate atomic diameter variations between Ni and some other elements and solubility data for these elements in Ni at 1000 ◦C. As can be seen, Ta, W, Mo, Ti, V, Cr, and Mn provide the best combination of atomic radii mismatch and appreciable solubility needed for solid solution strengthening. Therefore, Mo, Ti, Cr, and Mn are frequently used in many single phase Ni-base alloys for solid solution strengthening. Fe, Co and Cu also can be used for solid solution strengthening due to their high solubility in Ni [83]. The solid solution strengthening from additions of these elements will be effective only if the element concentrations do not exceed their solubility limits in Ni-based austenite, otherwise particle precipitation will take place. Although Ti, Al, and Nb can be effective solid solution strengtheners, they typically improve strength via precipitation of γ0-Ni3(Ti,Al) or Ni3Nb intermetallic particles, or carbides [2]. In addition to the atomic size misfit, electron configuration [84], diffusion coefficients [39], and γ-γ0 lattice misfit (dependent on element partitioning between γ-matrix and γ0-particles) [85] were shown to influence the solid solution strengthening at high temperatures. Thus, slow diffusing elements Re, Os, Ir, Tc, Ru, and W provide the best high temperature strength and creep resistance.

Table 2. Approximate atomic size variations and solubility limits of some elements in Ni.

Approximate Atomic Size Approximate Solubility in Ni at Element Difference Compared to Ni, % 1000 ◦C, wt.% [34] Ta 51 14 − W 44 38 − Mo 41 34 − Ti 27 10 − V 22 20 − Cr 17 40 − Mn 12 20 − Nb 49 6 − Al 31 7 Si 38 8 C 82 0.2 Fe 9 100 − Co 3 100 − Cu 4 100

4.3. Precipitation Strengthening In Monel K500 significant hardening is associated with the precipitation of gamma prime (γ’) particles. The γ’ precipitates have an ordered fcc crystal structure; therefore, the dislocation interaction mechanisms with the ordered precipitates dictate the level of strengthening. Initially, the particles nucleate and grow at a relatively short distance from each other. In this case of small interparticle spacing the particle shearing mechanism prevails. With coarsening, as the interparticle spacing increases, dislocation looping between the particles starts dominating. When precipitate shearing occurs, the strengthening of an alloy can arise from the (a) lattice mismatch between particles and the matrix, (b) order hardening, and (c) modulus mismatch. In Monel K500, the order hardening prevails over other shearing mechanisms. The lattice mismatch contribution to the total precipitation strengthening effect is small, because the coherency strain value is extremely small (less than 0.003) [69]. The precipitation strengthening mechanism may change with aging, due to the precipitate size variation with aging. In the under-aged and peak-aged condition, precipitation strengthening is governed predominantly by the particle shearing. This is associated with a lower shear stress required for particle shearing compared to the Orowan stress arising during looping of dislocations around precipitates [86]. Metals 2020, 10, 1358 8 of 18

In the overaged condition, two processes may operate simultaneously or one after another: (i) shearing in slightly overaged conditions and (ii) looping in significantly overaged conditions. When a dislocation pair approaches a particle surrounded by a loop, the first dislocation is blocked by this dislocation loop. The second dislocation approaching the first increases the shear stress between the first dislocation and the loop. Consequently, the particle shearing (accompanied by the loop annihilation) occurs even when the looping mechanism dominates. Thus, planar slip may be observed in an alloy containing slightly overaged γ’ particles. The reduction in stress due to the coarsening of precipitates is compensated by the transition of dislocation pairs to single dislocations, and thus, there is a net increase in strength in the initial stages of overaging. Order hardening originates from the interaction of dislocations with ordered precipitates (Figure8). The cutting of an ordered precipitate by the first dislocation shifts one half of a particle crystal plane by one interatomic distance against the other half plane (Figure8b). This results in creation of an antiphase boundary (APB) across the slip plane that represents an atomic layer of incorrect order. After the second dislocation passing through the same particle, the second shift of a crystal half plane occurs (Figure8c,d). This restores the atomic order. Thus, dislocations travelling in pairs (often referred to as superlattice dislocations) maintain the ordered structure after Metalsprecipitate 2020, 10 shearing., x FOR PEER REVIEW 9 of 19

Figure 8. Schematic illustration of dislocations interacting with an ordered Ni3Al precipitate [83]: (a) Figure 8. Schematic illustration of dislocations interacting with an ordered Ni3Al precipitate [83]: (a) the original condition prior to particle cutting, (b) the first dislocation penetrated a particle, (c) the the original condition prior to particle cutting, (b) the first dislocation penetrated a particle, (c) the second dislocation penetrated a particle, and (d) shifting of one half of the particle against the other second dislocation penetrated a particle, and (d) shifting of one half of the particle against the other half by two interatomic distances restored the atomic order . half by two interatomic distances restored the atomic order . Figure9 shows examples of such paired superlattice dislocations and particle cutting in a Ni-base alloy.Figure This deformation 9 shows examples involves of cross-slipsuch paired of segmentssuperlattice of thedislocations superlattice and dislocations particle cutting from thein a {111} Ni- baseslip plane alloy. to This {001} deformation cross-slip plane. involves These cross cross-slipped-slip of segments segments of the cannot superlattice move without dislocations forming from an APB the {111}and thus slip resistplane deformation.to {001} cross- Strengtheningslip plane. These from cross cross-slip-slipped becomes segments more cannot important move without with increasing forming antemperature APB and becausethus resist the deformation. cross-slip is thermally Strengthening activated. from cross-slip becomes more important with increasingThe initial temperature cutting of because an ordered the particlecross-slip by is a singlethermally dislocation activated. requires a stress increase. This stress increase provides energy for formation of the antiphase boundary. With increasing the antiphase boundary energy larger forces are required to cut the precipitates. Thus, the yield stress increases. The antiphase boundary energy provides a significant contribution to strengthening and was observed varying with alloy composition and particle chemistry. In particular, Ti, Co, Mo, and Fe were reported to effectively increase the antiphase boundary energy [87]. Finer particles (<20 nm) may be too small to hold two dislocations (Figure 10a), this is called a weak pair coupling. With larger particles, the second dislocation may start interacting with the particle before the first dislocation moves away (Figure 10b), this is called a strong pair coupling. In both cases, the strengthening contribution is due to the creation of an APB within the precipitate. The dislocation–particle interaction mechanism changes from shearing to looping with a particle size increase. Figure 10c qualitatively illustrates a dependence of precipitation strengthening on particle Figure 9. TEM images showing (a) paired super lattice dislocations and (b) γ′ precipitate being cut by the moving dislocations in a Ni superalloy [83].

The initial cutting of an ordered particle by a single dislocation requires a stress increase. This stress increase provides energy for formation of the antiphase boundary. With increasing the antiphase boundary energy larger forces are required to cut the precipitates. Thus, the yield stress increases. The antiphase boundary energy provides a significant contribution to strengthening and was observed varying with alloy composition and particle chemistry. In particular, Ti, Co, Mo, and Fe were reported to effectively increase the antiphase boundary energy [87]. Finer particles (<20 nm) may be too small to hold two dislocations (Figure 10a), this is called a weak pair coupling. With larger particles, the second dislocation may start interacting with the particle before the first dislocation moves away (Figure 10b), this is called a strong pair coupling. In both cases, the strengthening contribution is due to the creation of an APB within the precipitate. The dislocation–particle interaction mechanism changes from shearing to looping with a particle size increase. Figure 10c qualitatively illustrates a dependence of precipitation strengthening on particle size: this dependence shows a maximum. As the particle size varies with ageing, the precipitation strengthening contribution depends on the aging temperature and time [68]. At first, precipitation of the γ’ from the supersaturated Ni matrix leads to an increase in strength with increasing temperature and time. However, with further increase in temperature/time, an over-aging may lead to the particle Metals 2020, 10, x FOR PEER REVIEW 9 of 19

Metals 2020, 10, 1358 9 of 18

Figure 8. Schematic illustration of dislocations interacting with an ordered Ni3Al precipitate [83]: (a) the original condition prior to particle cutting, (b) the first dislocation penetrated a particle, (c) the size: this dependence shows a maximum. As the particle size varies with ageing, the precipitation second dislocation penetrated a particle, and (d) shifting of one half of the particle against the other strengthening contribution depends on the aging temperature and time [68]. At first, precipitation of half by two interatomic distances restored the atomic order . the γ’ from the supersaturated Ni matrix leads to an increase in strength with increasing temperature and time. However, with further increase in temperature/time, an over-aging may lead to the particle Figure 9 shows examples of such paired superlattice dislocations and particle cutting in a Ni- coarsening and a decrease in particle volume fraction (Figure 10d). A decrease in particle volume base alloy. This deformation involves cross-slip of segments of the superlattice dislocations from the fraction should be avoided as it results in a decrease in potential number of dislocation–particle {111} slip plane to {001} cross-slip plane. These cross-slipped segments cannot move without forming interaction sites and strengthening capacity. Therefore, a careful design of heat treatment schedule is an APB and thus resist deformation. Strengthening from cross-slip becomes more important with required to simultaneously maximise the particle size and particle volume fraction. increasing temperature because the cross-slip is thermally activated.

Metals 2020, 10, x FOR PEER REVIEW 10 of 19 coarsening and a decrease in particle volume fraction (Figure 10d). A decrease in particle volume fraction should be avoided as it results in a decrease in potential number of dislocation–particle interactionFigure 9.sites TEMTEM and images images strengthening showing showing ( a )capacity.) paired paired super super Therefore, lattice dislocations a careful design and (b) of γ′γ0 heatprecipitateprecipitate treatment being schedule cut by is requiredthe moving to simultaneously dislocations in maximise a Ni superalloy the particle [83]. size and particle volume fraction.

The initial cutting of an ordered particle by a single dislocation requires a stress increase. This stress increase provides energy for formation of the antiphase boundary. With increasing the antiphase boundary energy larger forces are required to cut the precipitates. Thus, the yield stress increases. The antiphase boundary energy provides a significant contribution to strengthening and was observed varying with alloy composition and particle chemistry. In particular, Ti, Co, Mo, and Fe were reported to effectively increase the antiphase boundary energy [87]. Finer particles (<20 nm) may be too small to hold two dislocations (Figure 10a), this is called a weak pair coupling. With larger particles, the second dislocation may start interacting with the particle before the first dislocation moves away (Figure 10b), this is called a strong pair coupling. In both cases, the strengthening contribution is due to the creation of an APB within the precipitate. The dislocation–particle interaction mechanism changes from shearing to looping with a particle size increase. Figure 10c qualitatively illustrates a dependence of precipitation strengthening on particle size: this dependence shows a maximum. As the particle size varies with ageing, the precipitation strengthening contribution depends on the aging temperature and time [68]. At first, precipitation of the γFigure’ from 10. 10.the Schematic Schematicsupersaturated illustration illustration Ni matrixof of(a) ( aweak) leads weak dislocation to dislocation an increase coupling, coupling, in strength (b) strong (b) strong with dislocation dislocationincreasing coupling coupling, temperature, and and time.(andc) strengthening ( cH) strengtheningowever, wit effecth e furtherff variationect variation increase with with particle in particle temperature size size [46 [];46 /(];time,d ()d e)ffect e anffect overof of aging aging-aging time time may on on leadthe the γ′γ to0 sizesize the and particle volume fraction in Monel K500 for various aging temperatures:temperatures: 1–6001–600 ◦°C,C, 2 2–650–650 °C,◦C, 3 3–700–700 °C◦C[ [8888].].

Precipitate morphology maymay changechange withwith aa γγ/γ′/γ0 lattice mismatch ( (FigureFigure 1111)[) [89].]. The The precipitate precipitate shape varies to minimise strain in the vicinity of particles and the particle surface energy.energy. At low levels of lattice mismatch, the strain is low and the spherical shape is favourable for the decrease of surface energy. At At higher levels of lattice mismatch, the strain increases and becomes orientation dependent. T Thehe orientation relationship between the γ--matrixmatrix and γγ′0--particleparticle promotes a cuboidal shape. These These γ′γ0--particlesparticles typical typicallyly align align along along the the <100<100>> directionsdirections of of the the matrix matrix which which have the lowest elastic elastic sti stiffness.ffness. The The precipitate precipitate morphology morphology also also varies varies with with coarsening. coarsen Thus,ing. Thus, on initial on stagesinitial stagesof nucleation of nucleation and growth and growth the particles the particles are spherical, are spherical, and then and gradually then gradually change to change cubes, toarrays cubes, of arrayscubes, of and cubes, dendritic and dendritic shape as shape coarsening as coarsening proceeds. proceeds. The lattice The mismatch lattice mismatch between between particles particles and the andmatrix the contributes matrix contributes to strengthening to strengthening of an alloy. of an alloy.

Figure 11. Optical images of γ′ particles in a Ni-Al-Mo alloy at different lattice mismatches overlaid in the top right corner [89]. Metals 2020, 10, x FOR PEER REVIEW 10 of 19 coarsening and a decrease in particle volume fraction (Figure 10d). A decrease in particle volume fraction should be avoided as it results in a decrease in potential number of dislocation–particle interaction sites and strengthening capacity. Therefore, a careful design of heat treatment schedule is required to simultaneously maximise the particle size and particle volume fraction.

Figure 10. Schematic illustration of (a) weak dislocation coupling, (b) strong dislocation coupling, and (c) strengthening effect variation with particle size [46]; (d) effect of aging time on the γ′ size and volume fraction in Monel K500 for various aging temperatures: 1–600 °C, 2–650 °C, 3–700 °C [88].

Precipitate morphology may change with a γ/γ′ lattice mismatch (Figure 11) [89]. The precipitate shape varies to minimise strain in the vicinity of particles and the particle surface energy. At low levels of lattice mismatch, the strain is low and the spherical shape is favourable for the decrease of surface energy. At higher levels of lattice mismatch, the strain increases and becomes orientation dependent. The orientation relationship between the γ-matrix and γ′-particle promotes a cuboidal shape. These γ′-particles typically align along the <100> directions of the matrix which have the lowest elastic stiffness. The precipitate morphology also varies with coarsening. Thus, on initial stages of nucleation and growth the particles are spherical, and then gradually change to cubes, arrays of cubes, and dendritic shape as coarsening proceeds. The lattice mismatch between particles Metals 2020, 10, 1358 10 of 18 and the matrix contributes to strengthening of an alloy.

Figure 11. OpticalFigure 11. imagesOptical imagesof γ′ particles of γ0 particles in in a aNi Ni-Al-Mo-Al-Mo alloy alloy at di ffaterent different lattice mismatches lattice mismatches overlaid in overlaid in the topMetals right the2020 top, 10corner, rightx FOR corner PEER[89]. REVIEW [89]. 11 of 19

TThehe γ′γ 0precipitateprecipitatess can can be be hardened hardened by by solid solid solution solution strengthening strengthening (Figure (Figure 12 12).). Mo, Mo, Si Si,, and and Ti Ti werewere shown to to be be the the most most eff ective effective elements elements for this. for Althoughthis. Although Cu and Cu Co and exhibit Co substantial exhibit substantial solubility solubilityin the particles, in the theyparticles do not, they increase do not the increase strength the of strengthγ0. The latticeof γ′. The parameter lattice parameter of both the ofγ matrixboth the and γ matrixγ0 precipitates and γ′ precipitates depends on depend type ands on concentrationstype and concentrations of alloying of elements.alloying elements. Various elementsVarious elements partition partitiondifferently differently to the matrix to the and matrix particles and andparticles change and their change lattice the parameters.ir lattice parameter Therefore,s. Therefore, the degree the of degreeγ/γ0 lattice of γ/γ′ mismatch lattice mismatch depends ondepends the relative on the partitioning relative partitioning of alloying of elements alloying and elements lattice parametersand lattice parametersvariations with variations alloying wi (Figureth alloying 13). (Figure 13).

Figure 12. a γ b Figure 12. (a() )Solubility Solubility of of various various elements elements in in γ′ 0particlesparticles at at 1150 1150 °C◦C in in Ni Ni-Al-X-Al-X alloy alloy system, system, (b ( ) )effect effect of various elements on the hardness of γ particles at 25 C[90]. of various elements on the hardness of γ′ 0particles at 25 °C◦ [90].

Figure 13. Effect of Ti/Al ratio on the γ/γ′ mismatch for Ni-20Cr-Al-Ti and Ni-20Cr-5Mo-Al-Ti alloys with a constant (Al + Ti) content of 3.5 wt.% [34].

4.4. Dislocation Strengthening (Work Hardening) Ni-Cu alloys exhibit rapid work hardening, characteristic for all Ni-base alloys due to their fcc crystal structure [91]. This allows to increase the room temperature hardness and strength applying large strains of cold deformation. However, to illuminate fracture development, (i) the number of deformation passes in a processing technology should be controlled, (ii) strain per pass must be reduced with each deformation stage, and (iii) intermediate annealing should be applied to restore ductility for further processing [57]. The rate of work hardening in the presence of precipitates was found to be higher than in a precipitate-free (solution-treated) matrix and increased with increasing precipitate size [92]. In alloys where the particles were small and shearing mechanism operated, the increase in work hardening rate was minor. The passage of dislocations through particles reduced their diameters, making the Metals 2020, 10, x FOR PEER REVIEW 11 of 19

The γ′ precipitates can be hardened by solid solution strengthening (Figure 12). Mo, Si, and Ti were shown to be the most effective elements for this. Although Cu and Co exhibit substantial solubility in the particles, they do not increase the strength of γ′. The lattice parameter of both the γ matrix and γ′ precipitates depends on type and concentrations of alloying elements. Various elements partition differently to the matrix and particles and change their lattice parameters. Therefore, the degree of γ/γ′ lattice mismatch depends on the relative partitioning of alloying elements and lattice parameters variations with alloying (Figure 13).

MetalsFigure2020, 10 12., 1358 (a) Solubility of various elements in γ′ particles at 1150 °C in Ni-Al-X alloy system, (b) effect11 of 18 of various elements on the hardness of γ′ particles at 25 °C [90].

Figure 13. EffectEffect of TiTi/Al/Al ratio on the γγ/γ′/γ0 mismatch for Ni Ni-20Cr-Al-Ti-20Cr-Al-Ti and Ni-20Cr-5Mo-Al-TiNi-20Cr-5Mo-Al-Ti alloys with a constant (Al + Ti) content of 3.5 wt.% [34].

4.4. Dislocation S Strengtheningtrengthening ( (WorkWork Hardening)Hardening) Ni-CuNi-Cu alloys exhibit rapid work hardening, characteristic for all Ni-baseNi-base alloys due to their fcc crystal structure [91]. This allows to increase the room temperature hardness and strength applying large strains of coldcold deformation.deformation. However, to illuminate fracture development, (i) the number of deformation passes passes in in a processing a processing technology technology should should be controlled, be controlled, (ii) strain (ii) perstrain pass per must pass be must reduced be withreduced each with deformation each deformation stage, and stage, (iii) intermediate and (iii) intermediate annealing shouldannealing be applied should tobe restore applied ductility to restore for furtherductility processing for further [57 processing]. [57]. The rate of work hardening in the presence of precipitates was found to be higher than in a precipitate-freeprecipitate-free (solution-treated)(solution-treated) matrixmatrix andand increasedincreased withwith increasingincreasing precipitateprecipitate sizesize [[9292].]. In alloys where the particlesparticles were smallsmall andand shearingshearing mechanismmechanism operated,operated, the increase in work hardening rate was minor.minor. The passage of dislocations through particles reduced their diameters, making the passage of the following dislocations easier. A relative reduction in diameter for a small particle would be larger than that for a large particle. With growing particle size the work hardening rate increased, even when particle shearing occurred. Looping increases the work hardening rate. This can be explained by a gradual reduction in the interparticle spacing with the formation of dislocation loops around precipitates, making passage of the successive dislocations difficult [93]. With an increase in cold deformation the strength of Monel K500 increases and ductility decreases (Figure 14a). The optimum deformation strain for Monel K500 was found to be about 20% [88]: at this strain an increment to the yield stress is quite substantial (about 300 MPa) and the remaining ductility is still significant (about 30% of elongation to failure). Cold deformation has been shown to affect the strength variation during aging (Figure 14b). With an increase in aging temperature up to 600 ◦C, the strength increased for 0 and 10% cold worked samples. However, the samples cold worked to 20% and 25% showed strength peaks in the range of 560–600 ◦C followed by the strength decrease. Obviously with an increase in cold deformation the strain induced precipitation occurs faster; this leads to faster particle growth accompanied by a decrease in particle number density and strength decrease. Irrespective of the aging temperature the strength showed a minimum for 10% of prior cold worked material (Figure 14c). This could be explained if small amounts of deformation facilitated the dislocation breaking from their pinning points, increasing the dislocation mobility and decreasing the yield stress. Since work hardening is related to solute element concentrations and presence of precipitates, the work hardening rate generally increases with complexity of alloy composition. Similarly, the age-hardenable alloys have higher work hardening rates than their solid solution equivalents [94]. In the annealed condition, Monel K500 can be cold-worked using traditional processing technology. It has excellent ductility, although deformation forces may be high. Work hardening of Monel K500 Metals 2020, 10, x FOR PEER REVIEW 12 of 19

passage of the following dislocations easier. A relative reduction in diameter for a small particle would be larger than that for a large particle. With growing particle size the work hardening rate increased, even when particle shearing occurred. Looping increases the work hardening rate. This can be explained by a gradual reduction in the interparticle spacing with the formation of dislocation loops around precipitates, making passage of the successive dislocations difficult [93]. With an increase in cold deformation the strength of Monel K500 increases and ductility decreases (Figure 14a). The optimum deformation strain for Monel K500 was found to be about 20% [88]: at this strain an increment to the yield stress is quite substantial (about 300 MPa) and the remaining ductility is still significant (about 30% of elongation to failure). Cold deformation has been shown to affect the strength variation during aging (Figure 14b). With an increase in aging temperature up to 600 °C, the strength increased for 0 and 10% cold worked samples. However, the samples cold worked to 20% and 25% showed strength peaks in the range of 560–600 °C followed by the strength decrease. Obviously with an increase in cold deformation the strain induced Metals 2020, 10, 1358 12 of 18 precipitation occurs faster; this leads to faster particle growth accompanied by a decrease in particle number density and strength decrease. Irrespective of the aging temperature the strength showed a withminimum cold deformationfor 10% of prior is shown cold workein Figured material 15 in comparison (Figure 14c with). This other could materials: be explained Monel if K500, small precipitationamounts of deformation hardenable facilitated with Ni3(Ti,Al) the dislocation and TiC, approachesbreaking from Inconel their pinning 600, precipitation points, increasing hardenable the withdislocation Cr-rich mobility carbides. and decreasing the yield stress.

a b c

Figure 14. (a) Mechanical properties variation with cold deformation, (b) effect of age hardening Figure 14. (a) Mechanical properties variation with cold deformation, (b) effect of age hardening temperatures on properties for various strains of prior cold deformation: 1–0, 2–0.10, 3–0.20, and 4–0.25, temperatures on properties for various strains of prior cold deformation: 1–0, 2–0.10, 3–0.20, and 4– and (c) effect of prior cold deformation on the yield stress after aging at various temperatures, all for 0.25, and (c) effect of prior cold deformation on the yield stress after aging at various temperatures, MetalsMonel 2020, 10 K500, x FOR [88 PEER]. REVIEW 13 of 19 all for Monel K500 [88].

Since work hardening is related to solute element concentrations and presence of precipitates, the work hardening rate generally increases with complexity of alloy composition. Similarly, the age- hardenable alloys have higher work hardening rates than their solid solution equivalents [94]. In the annealed condition, Monel K500 can be cold-worked using traditional processing technology. It has excellent ductility, although deformation forces may be high. Work hardening of Monel K500 with cold deformation is shown in Figure 15 in comparison with other materials: Monel K500, precipitation hardenable with Ni3(Ti,Al) and TiC, approaches Inconel 600, precipitation hardenable with Cr-rich carbides.

FigureFigure 15. 15.Comparison Comparison of of workwork hardeninghardening capacity of of Monel Monel K500 K500 to toother other Ni- Ni-basebase alloys alloys [53]. [53 ] (reproduced with a permission from the Special Metals Corporation). 5. Mechanical Properties 5. Mechanical Properties Room temperature tensile properties of Monel 400 and K500 are shown in Table 3 [20]. In hot- rolledRoom and temperature annealed conditions tensile properties Monel ofK500 Monel shows 400 and a 10 K5000–200 are MPa shown higher in Table yield3[ 20 stress]. In hot-rolled (YS) and andultimate annealed tensile conditions strength Monel(UTS) than K500 Monel shows 400 a 100–200, due to MPa solid higher solution yield and stress precipitation (YS) and ultimatestrengthening. tensile strengthIn both alloys (UTS) YS than and Monel UTS 400, are dueslightly to solid higher solution for cold and fo precipitationrmed products strengthening., due to work In bothhardening alloys. YSIn andhot UTS finished are slightlyMonel higher K500, for annealing cold formed may products, lead to duestrength to work decrease hardening. by about In hot 30 finished%, following Monel K500,dissolution annealing of precipitates may lead toand strength dislocation decrease annihilation. by about 30%,In contrast, following age dissolution-hardening ofmay precipitates result in 1. and3– dislocation2.5 times increase annihilation. in strength In contrast, due to age-hardeningprecipitation. For may the result cold in formed 1.3–2.5 products, times increase annealing in strength decreases due tostrength precipitation. by about For 40 the–50 cold%; although formed, products, age-hardening annealing may decreases increase strength by aboutup to 1.3 40–50%; times. although, As seen, age-hardeningthe age-hardening may heat increase treatment strength is more by up effective to 1.3 times. in increasing As seen, strength the age-hardening of the hot finished heat treatment products, is morecompared effective to the in cold increasing finished. strength of the hot finished products, compared to the cold finished. High and low temperature tensile properties of Monel K500 are shown in Figure 16. For hot rolled product, with an increaseTable in temperature 3. Mechanical the propert YS andies UTS of Monel do not 400 vary and significantlyK500 [20]. until 650 ◦C (1200 ◦F) and 150 C (300 F), respectively. Age-hardening increases YS and UTS by about 1.5 and 2 times, ◦ ◦ Yield Strength, Tensile Elongation, respectively.Processing However, Condition the YS stability at high temperature decreases, YS starts going downHB at a lower MPa Strength, MPa % temperature of 430 ◦C (800 ◦F) compared to the non-aged material. In contrast, the stability of UTS Monel 400 Hot-Finished 280–690 550–760 30–60 140–240 Hot-Finished, 170–340 520–620 35–60 110–150 Annealed Cold-Drawn 380–690 580–830 22–40 160–225 Monel K500 Hot-Finished 280–760 620–1070 45–20 140–315 Hot-Finished, Aged 690–1034 965–1310 30–20 265–346 Hot-Finished, 280–414 621–760 45–25 140–185 Annealed Hot-Finished, Annealed 586–830 896–1140 35–20 250–315 and Aged Cold-Drawn, As- 483–860 690–965 35–13 175–260 Drawn Cold-Drawn, Aged 655–1100 931–1280 30–15 255–370 Cold-Drawn, 280–414 621–760 50–25 140–185 Annealed Cold-Drawn, 586–830 896–1310 30–20 250–315 Annealed and Aged

High and low temperature tensile properties of Monel K500 are shown in Figures 16. For hot rolled product, with an increase in temperature the YS and UTS do not vary significantly until 650 °C Metals 2020, 10, 1358 13 of 18

increases, UTS starts decreasing at a higher temperature of 300 ◦C (550 ◦F) compared to the non-aged material. The decrease in YS stability after ageing can be associated with destruction of the dislocation sub-structure during ageing, and the increase in UTS stability follows particle precipitation. Annealing prior age-hardening does not provide any strength gain compared to the hot-rolled and age-hardened product, however the stability of elongation increase. This is a consequence of significantly reduced dislocation density after annealing and increased distance of dislocations free pass. With a decrease in temperature, the tensile strength and yield stress both increase while ductility and toughness remain virtually the same (Figure 16d). No ductile-to-brittle transition occurs even at temperatures as low as Metalsthat of2020 liquid, 10, x hydrogen.FOR PEER REVIEW Thus, this alloy is suitable for many cryogenic applications. 14 of 19

(1200 °F) and 150 °C (300Table °F), 3.respectively.Mechanical properties Age-hardening of Monel increases 400 and K500YS and [20]. UTS by about 1.5 and 2 times, respectively. However, the YS stability at high temperature decreases, YS starts going down at Processing Condition Yield Strength, MPa Tensile Strength, MPa Elongation, % HB a lower temperature of 430 °C (800 °F) compared to the non-aged material. In contrast, the stability Monel 400 of UTS increases,Hot-Finished UTS starts decreasing 280–690 at a higher temperature 550–760 of 300 °C (550 °F) 30–60 compared 140–240 to the non-agedHot-Finished, material. Annealed The decrease in YS 170–340 stability after ageing can 520–620 be associated with 35–60 destruction 110–150 of the dislocationCold-Drawn sub-structure during ageing 380–690, and the increase 580–830 in UTS stability 22–40 follows 160–225particle Monel K500 precipitation.Hot-Finished Annealing prior age-hardening 280–760 does not provide 620–1070 any strength gain 45–20 compared 140–315 to the hot-rolledHot-Finished, and age Aged-hardened product, 690–1034 however the stability 965–1310 of elongation increase. 30–20 This 265–346 is a consequenceHot-Finished, of significantly Annealed reduced 280–414dislocation density after 621–760 annealing and increased 45–25 distance 140–185 of Hot-Finished, Annealed and Aged 586–830 896–1140 35–20 250–315 dislocationsCold-Drawn, free As-Drawn pass. With a decrease 483–860 in temperature, the 690–965 tensile strength and 35–13 yield stress 175–260 both increaseCold-Drawn, while ductility Aged and toughness 655–1100 remain virtually the 931–1280same (Figure 16d). No 30–15 ductile-to 255–370-brittle transitionCold-Drawn, occurs Annealedeven at temperatures 280–414 as low as that of liquid 621–760 hydrogen. Thus, this 50–25 alloy is suitable 140–185 Cold-Drawn, Annealed and Aged 586–830 896–1310 30–20 250–315 for many cryogenic applications.

a b c d

Figure 16. High temperature tensile properties of Monel K500: (a) hot-rolled, (b) hot-rolled and Figure 16. High temperature tensile properties of Monel K500: (a) hot-rolled, (b) hot-rolled and age- age-hardened, and (c) annealed and age-hardened; (d) low temperature tensile properties of Monel hardened, and (c) annealed and age-hardened; (d) low temperature tensile properties of Monel K500 K500 [53] (reproduced with a permission from the Special Metals Corporation). [53]. In view of excellent corrosion resistance coupled with high strength and toughness at cryogenic In view of excellent corrosion resistance coupled with high strength and toughness at cryogenic temperatures [95,96] Monel alloys are good candidates for structural components of machinery and temperatures [95,96] Monel alloys are good candidates for structural components of machinery and storage of liquefied gases in aerospace and chemical industry. High temperature application of Monel storage of liquefied gases in aerospace and chemical industry. High temperature application of Monel is limited by low melting temperature of Cu. However, the precipitation strengthening capacity in the is limited by low melting temperature of Cu. However, the precipitation strengthening capacity in Ni-Cu alloy system can be improved with appropriate alloying element additions and heat treatment. the Ni-Cu alloy system can be improved with appropriate alloying element additions and heat Ni-Cu alloys are easily weldable [97] and were successfully used as an input material in powder treatment. Ni-Cu alloys are easily weldable [97] and were successfully used as an input material in based [98] and wire arc additive manufacturing [99]. Development of these modern technologies powder based [98] and wire arc additive manufacturing [99]. Development of these modern allows to apply the Ni-Cu alloys as a surface cladding material for protection of less corrosion resistant technologies allows to apply the Ni-Cu alloys as a surface cladding material for protection of less core or in components with mechanical properties gradient. corrosion resistant core or in components with mechanical properties gradient.

6. Conclusions Due to Ni and Cu exhibiting complete solid solubility, their alloys are single phase. Thus, in Ni- Cu alloy four strengthening mechanisms operate: grain refinement, solid solution, precipitation, and dislocation strengthening (work hardening). Composition of an alloy determines whether the solid solution or precipitation strengthening dominates. Mo, Ti, Cr, and Mn are the most frequently used solid solution strengthening elements, due to their significant difference in atomic sizes from Ni. Such elements as Fe, Co, and Cu are the second order solid solution strengtheners, due to their high solubility in Ni. Al and Ti are the most effective precipitation strengthening elements in Ni-Cu alloys, as they tend to form NiAlTi-rich intermetallic particles. Sometimes Mn-rich M23C6 or Ti-rich MC carbides can also precipitate. The dislocation strengthening capacity of Ni-Cu alloys is substantial, which is determined by their fcc crystal structure. However, cold deformation above 20% would decrease ductility below the practically reasonable limits of 30% elongation. In presence of particle Metals 2020, 10, 1358 14 of 18

6. Conclusions Due to Ni and Cu exhibiting complete solid solubility, their alloys are single phase. Thus, in Ni-Cu alloy four strengthening mechanisms operate: grain refinement, solid solution, precipitation, and dislocation strengthening (work hardening). Composition of an alloy determines whether the solid solution or precipitation strengthening dominates. Mo, Ti, Cr, and Mn are the most frequently used solid solution strengthening elements, due to their significant difference in atomic sizes from Ni. Such elements as Fe, Co, and Cu are the second order solid solution strengtheners, due to their high solubility in Ni. Al and Ti are the most effective precipitation strengthening elements in Ni-Cu alloys, as they tend to form NiAlTi-rich intermetallic particles. Sometimes Mn-rich M23C6 or Ti-rich MC carbides can also precipitate. The dislocation strengthening capacity of Ni-Cu alloys is substantial, which is determined by their fcc crystal structure. However, cold deformation above 20% would decrease ductility below the practically reasonable limits of 30% elongation. In presence of particle forming elements in alloy composition, age hardening heat treatment is frequently used as the final strengthening operation.

Author Contributions: O.M. carried out the literature analysis and paper writing; A.K. contributed to the discussion of results. All authors have read and agreed to the published version of the manuscript. Funding: This research received no external funding. Acknowledgments: Olexandra Marenych is thankful for a scholarship sponsored by the Defence Materials Technology Centre at the University of Wollongong (Australia), which was established under the initiative of the Australian Government’s Defence Future Capability Technology Centre. Conflicts of Interest: The authors declare no conflict of interest.

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