Effects of Composition, Processing and Structure on ...

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Oct. 21, 2024

Effects of Composition, Processing and Structure on ...

Elemental nickel is used principally as an alloying element to increase the corrosion resistance of commercial iron and copper alloys; only about 13% of annual consumption is used in nickel-base alloys. Approximately 60% is used in stainless steel production, with another 10% in alloy steels and 2.5% in copper alloys. Nickel is also used in special-purpose alloys: controlled expansion, electrical resistance, magnetic, and shape memory alloys.

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Nickel and nickel alloys are used in the chemical processing, pollution control, power generation, electronics, and aerospace industries, which are taking advantage of their excellent corrosion, oxidation, and heat resistance.

Nickel is ductile and can be made by the conventional processing methods into cast, P/M, and various wrought products: bar/wire, plate/sheet, and tube. Commercially pure nickel has moderately high values of melting temperature (°C), density (8.902 g/cm3), and elastic modulus (204 GPa). It is ferromagnetic, with a Curie temperature of 358°C (676°F) and good electrical (25% IACS) and thermal conductivity (82.9 W/m K, or 48 Btu/ft h °F).

Effects of Alloying Elements in Nickel Alloys

Nickel has an face-centered cubic crystal (fcc) structure, to which it owes its excellent ductility and toughness. Because nickel has extensive solid solubility for many alloying elements, the microstructure of nickel alloys consists of the fcc solid-solution austenite (γ) in which precipitate particles can form.

Nickel forms a complete solid solution with copper and has nearly complete solubility with iron. It can dissolve about 35% Cr, about 20% each of molybdenum and tungsten, and about 5 to 10% each of aluminum, titanium, manganese, and vanadium. Thus, the tough, ductile fcc matrix can dissolve extensive amounts of elements in various combinations to provide solution hardening as well as improved corrosion and oxidation resistance. The degree of solution hardening has been related to the atomic size difference between nickel and the alloying element, and therefore the ability of the solute to interfere with dislocation motion.

Tungsten, molybdenum, niobium, tantalum, and aluminum, when aluminum is left in solution, are strong solution hardeners, with tungsten, niobium, tantalum, and molybdenum also being effective at temperatures above 0.6 T m (T m = melting temperature), where diffusion-controlled creep strength is important. Iron, cobalt, titanium, chromium, and vanadium are weaker solution-hardening elements. Aluminum and titanium are usually added together to form the age-hardening precipitate, Ni3(Al, Ti).

In addition, some alloying elements can partition to γ&#;, affecting the interface mismatch and precipitate-coarsening kinetics as well as contributing a solution-hardening component to strength, with titanium being the most effective at room and elevated temperatures.

However, titanium, niobium, and tantalum can influence mechanical properties still further by encouraging the formation of other similar types of precipitates. With higher titanium content, γ&#; will transform to the hexagonal close-packed (hcp) η- phase, Ni3Ti, which has an acicular or cellular morphology. With increased amounts of niobium, γ&#; transforms to the commercially important metastable body-centered tetragonal (bct) phase γ". A decrease in hardening will result if the equilibrium orthorhombic phase, Ni3Nb, is allowed to form. The actual phases precipitated and their effectiveness in hardening the micro-structure are dependent on the alloy composition, the applied heat treatments, the resulting precipitate volume fraction, and the service conditions.

Carbides. Although not a carbide former, nickel dissolves many elements that readily form the carbides seen in nickel alloys (MC, M6C, M7C3, M23C6). The MC carbides (where M = W, Ta, Ti, Mo, Nb) are usually large, blocky, and undesirable. The M6C carbides (M = Mo, W) can precipitate as small platelets in the grains or as blocky particles in boundaries useful for grain control, but deleterious for ductility and stress rupture properties. The M7C3 (M = Cr) can be useful when precipitated as discrete particles, but more so are grain boundary particles of M23C6 (M = Cr, Mo, W), where they can enhance creep rupture properties.

If carbides are allowed to agglomerate or form grain-boundary films during heat treatment or in service at elevated temperatures, they can seriously impair ductility and cause embrittlement. As in stainless steels, precipitation of chromium carbides at boundaries can lead to intergranular corrosion due to the chromium-depleted zone alongside the grain boundary becoming anodic to the rest of the grains.

This grain-boundary sensitization is controlled in several ways:

by avoiding the chromium-carbide aging temperature range (425 to 760°C) during processing,
with stabilization heat treatments to tie up carbon with more stable carbide formers (niobium, tantalum, titanium), and
by reducing the carbon level in the base alloy.

Nickel alloys

Nickel is alloyed to extend the good corrosion resistance and good heat resistance of elemental nickel. Even with extensive amounts of alloying elements, the tough, ductile fcc austenitic matrix is preserved.

It is convenient to describe nickel alloys by grouping them into their two broad application areas: corrosion resistance, especially in aqueous environments, and heat resistance. Naturally, this artificial separation should not be considered a rigid barrier as the corrosion-resistant alloys have good strength above room temperature and the heat-resistant alloys have good corrosion resistance. The unique, special-property alloys, many of which are also used for their good corrosion and heat resistance as well as high strength, are described separately.

Corrosion-Resistant Nickel Alloys. The commercially pure nickel grades, Nickel 200 to 205, are highly resistant to many corrosive media, especially in reducing environments, but also in oxidizing environments where they can maintain the passive nickel oxide surface film. They are used in the chemical processing and electronics industries.

They are hot worked at 650 to °C, annealed at 700 to 925 °C, and are hardened by cold working. For processed sheet, for example, the tensile properties in the annealed condition (460 MPa, tensile strength; 148 MPa, yield strength; and 47% elongation) can be increased by cold rolling up to 760 MPa tensile strength, 635 MPa yield strength, and 8% elongation.

Because of its nominal 0.08% C content (0.15% max), Nickel alloy 200 (UNS No ) should not be used above 315°C, since embritlement results from the precipitation of graphite in the temperature range 425 to 650°C. Higher-purity nickel is commercially available for various electrical applications.

The low-alloy nickels. These alloys contain 94% min Ni. The 5% Mn solid-solution addition in Nickel 211 protects against sulfur in service environments. As little as 0.005% S can cause liquid embrittlement at unalloyed nickel grain boundaries in the range between 640 and 740°C.

Duranickel, alloy 301 (Ni-4.5Al-0.6Ti), offers the corrosion resistance of commercially pure nickel with the strengthening provided by the precipitation of γ&#;. There is sufficient alloying additions in alloy 301 to lower the Curie temperature, making the alloy weakly ferromagnetic at room temperature.

The nickel-copper alloys are strong and tough, offering corrosion resistance in various environments, including brine and sulfuric and other acids, and showing immunity to chloride-ion stress corrosion. They are used in chemical processing and pollution control equipment. Capable of precipitating γ&#;, Ni3 (Al, Ti), with its 2.7Al - 0.6Ti alloy addition, alloy K-500 adds an age-hardening component to the good solution strengthening and work-hardening characteristics already available with the nominal 30% Cu in alloy 400. The composition of these alloys can be adjusted to decrease the Curie temperature to below room temperature.

The Ni-Cr-Fe (-Mo) alloys might simply be thought of as nickel-base analogs of the iron-base austenitic stainless steel alloys, with an interchange of the iron and nickel contents. In these commercially important alloys the chromium content in general ranges from 14 to 30% and iron from 3 to 20%. With a well-maintained Cr2O3 surface film, these alloys offer excellent corrosion resistance in many severe environments, showing immunity to chloride-ion stress-corrosion cracking. They also offer good oxidation and sulfidation resistance with good strength at elevated temperatures. These nickel-rich Ni-Cr-Fe alloys have maximum operating temperatures in the neighborhood of °C.

The Ni-Cr-(Fe)-Mo alloys consist of a large family of alloys that are used in the chemical processing, pollution control, and waste treatment industries to utilize their excellent heat and corrosion resistance. Alloys in this commercially important family, such as C-276 and alloy 625, are made even more versatile by their excellent welding characteristics and the corrosion resistance of welded structures.

The molybdenum additions to these alloys improve resistance to pitting and crevice corrosion. Aluminum improves the protective surface oxide film, and the carbide formers titanium and niobium are used to stabilize the alloys against chromium-carbide sensitization. Even with the low-level additions of aluminum and titanium to alloy 800, for example, small amounts of γ&#; can form in service during exposure to elevated temperatures. The high molybdenum and silicon additions in Hastelloy B and D promote good corrosion resistance during in the presence of hydrochloric and sulfuric acids.

Heat-Resistant Nickel Alloys. These nickel-containing materials include nickel-, iron-nickel-, or cobalt-base alloys. They can be made by wrought and P/M methods, and also with castings produced with carefully controlled conditions to provide the desired polycrystal, or elongated (directionally solidified), or single-crystal grain structure for improved elevated-temperature mechanical properties. The majority of the nickel-base superalloys utilize the combined strengthening of a solution-hardened austenite matrix with γ&#; precipitation.

The iron-base Fe-Ni-Cr heat-resistant alloys are extensions of the iron-base stainless steels with higher nickel and additions of other alloying elements. Retaining the fcc iron-nickel austenite matrix, these alloys (alloys A-286 and 901, for example) are workable into various wrought forms and are capable of precipitation hardening with γ&#;.

Alloys 903 and 909 are controlled thermal expansion Fe-Ni-Co-base alloys that are capable of age hardening with Ni3(Nb, Ti) precipitation and are designed to have high strength and low coefficient of thermal expansion for applications in gas turbine rings and seals up to 650°C.

These alloys are hot worked at about 870 to °C and solution heat treated at 815 to 980 °C. The standard aging treatment consists of 720 °C for 8 h, furnace cool at 55°C/h to 620°C for 8 h, followed by air-cooling. Alloy 909 in the as-hardened condition, for example, retains much of its room-temperature yield strength ( MPa) at 540°C, namely, 895 MPa.

Specialty Nickel Alloys. Unique combinations of properties are available with other nickel-base alloys for special applications. While some of these properties are also available to some extent with alloys described above, the alloys described below were developed to promote their rather unique properties.

There are many electrical resistance alloys used for resistance heating elements. They can contain 35 to 85% Ni, but invariably contain greater than 15% Cr to form an adherent surface oxide to protect against oxidation and carburization at temperatures up to to °C in air.

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Nickel Alloying in Carbon Steel: Fundamentals and ...

In the high-temperature range, the Fe&#;Ni binary alloy comprises a peritectic reaction such as the Fe&#;C binary alloy. Solidification from the liquid phase can proceed along various phase transformation scenarios ( Figure 1 c) [ 6 ]. For lower nickel additions (<3.4 wt.%), with a respectively low carbon content, solidification occurs entirely as-ferrite, subsequently transforming to austenite. The final transformation starts by nucleation of α-ferrite at the austenite grain boundaries and subsequent growth. The size of these ferrite grains depends on the degree of undercooling controlled by the Ni content and the cooling rate. At higher nickel content the austenite phase either transforms by nucleation of grain boundary ferrite or by diffusionless transformation into martensite below the Mtemperature. In the latter case, the boundaries of the prior columnar austenite grains will be maintained with the martensite substructure developing inside this boundary. Especially after direct solidification from liquid to austenite, impurities such as sulfur and phosphorous segregated to the austenite grain boundary also decorate the prior austenite grain boundary in the final martensitic microstructure. This is not the case when ferrite nucleation and growth erase the prior austenite boundaries in steels with lower nickel content. Additionally, dilatational stresses originating from the martensite transformation act on the prior austenite grain boundaries. The resulting embrittlement of such boundaries can cause intergranular cracking and is relevant to the performance of higher-nickel alloyed steels in as-cast condition as well as after fusion welding. The removal of this detrimental effect requires dedicated heat treatment.

The activation energy for the diffusion of nickel atoms in iron is rather high, so that the diffusion coefficient of nickel is generally low and comparable to the self-diffusion of iron [ 5 ]. Therefore, equilibration of concentration variations can be achieved only during long holding times at high temperature. However, nickel diffuses faster in bcc iron than in fcc iron. This difference causes a hysteresis in the transformation behavior between cooling and heating cycles and becomes pronounced for Ni contents above approximately 5%. Martensite transformation commences during down-cooling when passing the M(martensite-start) temperature. Before reaching M(martensite-finish) temperature, bcc and fcc phases coexist; however, the nickel concentration in either phase can be assumed to be nearly equal due to the very slow diffusion. However, when reheating above the(austenite-start) temperature into the region of coexisting fcc and bcc phases, nickel can partition into austenite, provided the holding time is sufficiently long. Such nickel-enriched crystallites can finally survive as metastable retained austenite phases when cooling back to ambient temperature. In technical steel alloys, carbon and manganese also partition to these austenite islands providing additional stabilization. The retained austenite fraction present in carbon steels after heat treatment enables a transformation-induced plasticity (TRIP) effect and makes an important contribution to enhanced formability and low temperature toughness, as will be discussed later (see Section 2.6 ).

Due to the similarity in principal properties, nickel and iron form a substitutional solid solution when mixed together. The Fe&#;Ni binary-phase diagram, which principally represents the possible phase formation in steels alloyed with nickel indicates the stabilization of the austenitic (fcc) phase in Ni-alloyed steel towards lower temperatures ( Figure 1 a). Two-phase ferritic-austenitic microstructures exist at room temperature when the Ni contents exceed approximately 15 wt.%. Bainite or martensite formation is promoted upon down-cooling as a result of the decreasing transformation temperature with increasing nickel content. Above 6 wt.% Ni content, diffusion-controlled transformation processes are fully suppressed, resulting in martensite transformation. In the presence of carbon, phase fields are somewhat shifted but, principally, the effects of nickel on the transformation behavior remain similar. Experimentally determined Aand Atemperatures are shown in Figure 1 b as a function of the nickel content in a 0.1 wt.% carbon-containing steel [ 4 ]. The influence of nickel is obviously stronger on the Atemperature than on the Atemperature. These temperatures are important with regard to re-austenitizing and tempering treatments.

Like silicon, nickel raises the activity of carbon in steel, which also affects the solubility of micro-alloy carbides at low level. The results of Eckstein et al. [ 15 ] indicate that the dissolution temperature of niobium carbo-nitride is increased by only about 5 K when adding 0.5 wt.%Ni to carbon steel. Manganese, molybdenum, and chromium on the contrary reduce the niobium carbonitride dissolution temperature by a similar magnitude at 0.5 wt.% addition levels. Sharma et al. [ 16 ] derived interaction parameters for nickel and other common alloying elements allowing the correction of the niobium solubility product. However, as the magnitude of the nickel effect on the solubility of micro-alloy carbides in relevant steel compositions is rather small, it has no significant consequences regarding practical soaking treatments. On the other hand, an enhanced carbon activity by nickel could promote micro-alloy precipitation at lower austenite temperature.

Nickel can form a carbide species with the composition NiC. The decomposition of the metastable NiC carbide into Ni and C starts at a temperature of around 465 °C and is accompanied by a thermal effect &#;ΔH &#; 10 kJ/mol [ 10 ]. However, due to its low affinity to carbon, nickel carbide does not form in steel. In ternary Fe&#;C&#;Ni alloys, the diffusivities of nickel and carbon are mutually enhanced. Small amounts of nickel dissolved in iron carbide (cementite) facilitate its decomposition, which results in graphite formation during long holding periods at high pearlite temperature. The graphite-promoting effect of nickel is prominent in cast iron alloys but not relevant to carbon steels. As early as the year it had been reported that nickel lowers the pearlite formation temperature and lowers the eutectoid carbon content [ 11 ]. Quantitatively, the addition of 1 wt.% Ni reduces the pearlite formation temperature by about 15 K and lowers the eutectoid carbon content by around 0.04 wt.% [ 12 ]. However, nickel does not refine the pearlite lamellar structure beyond that provoked by the undercooling effect [ 13 ]. Accordingly, nickel alloying to low carbon steel increases the phase fraction of pearlite, as was observed for a 0.17 wt.%C steel in which the pearlite phase fraction linearly increased from 0.2 to 0.3 when adding up to 3 wt.% Ni [ 14 ].

The nickel-induced lowering of the transformation temperature and the associated increase in the driving force for the bainitic transformation results in a refinement of the bainite package and block size, which noticeably supports the hardness-increasing effect of nickel [ 24 25 ]. Lowering the carbon content and increasing the nickel addition promotes this refinement. The resulting enhanced density of block boundaries and Bain variant pairs leads to a high hardness even at the low cooling rates encountered in the center section of heavy-gauged products. Alloying 2 wt.% Ni to a water-quenched high-strength steel allows the through-hardenable plate thickness to be more than tripled, as demonstrated by Figure 5

Nickel and manganese, extending the austenite phase field, act by mechanism (a). The solute drag mechanism (b) is most pronounced for large alloy atoms such as niobium, molybdenum, and titanium. The obstruction of carbon diffusion (c), retarding pearlite formation, is promoted by chromium and also by other carbide formers (Nb, Ti, V, Mo). Solute boron strongly segregates to the austenite grain boundary, lowering its energy, thus, obstructing ferrite nucleation by mechanism (d). The magnitude of mechanism (b) is influenced by the diffusivity of the respective drag exerting solutes at the temperature of phase transformation. Since nickel lowers the A r3 temperature, irrespective of the cooling rate, the diffusivity of these elements is inherently reduced. Regarding mechanism (d), the segregation of boron at the austenite grain boundary relies on the vacancy flux dragging boron (and also Mo, Nb) towards the boundary. As the austenite temperature decreases, vacancies flow from the grain interior to the boundary which acts as a sink. Again, the lower transformation temperature caused by nickel alloying enhances this non-equilibrium segregation mechanism by a stronger vacancy flux. Accordingly, for production routes where sufficiently high cooling rates can be guaranteed across the entire thickness, nickel alloying may not be essential for hardenability. In heavy-gauged products, however, the cooling rate in the center region is naturally limited and nickel alloying on top of the applicable maximum manganese content becomes relevant for retarding the transformation.

The multiplying factor of nickel quantifying its hardenability effect is moderate and comparable to that of silicon but smaller than those of manganese, molybdenum, and chromium [ 12 ]. However, nickel alloying at reasonably low addition levels can significantly enhance the efficiency of other alloying elements for generating strong microstructures by enabling synergy effects. To understand the principle of these synergies, the different mechanisms providing hardenability must be considered:

The martensite-start (M) temperature decreases with 17&#;19 K per weight percent nickel added [ 20 ]. It is commonly understood that martensite transformation is an athermal process (i.e., it does not require thermal activation) that cannot be suppressed. However, it has been known for a long time that isothermal formation of martensite is also feasible. Borgenstam and Hillert [ 21 ] have reviewed these early studies regarding the kinetics of isothermal martensite transformation. The activation energy was estimated to be in the range of 7 to 80 kJ/mol for different alloy systems. While activation energies on the lower side should rather be related to growth mechanisms involving dislocation movements, those on the higher side would quite well agree with the activation energy required for carbon diffusing away from a growing bcc plate within the undercooled austenite. The practical relevance of isothermal martensite formation has been concisely discussed by Villa and Somers [ 22 ] in the light of innovative cryogenic treatments on high-performance steels. Their results, using 15&#;17 wt.% Cr steel alloys, reveal that increasing the nickel content from 2 to 7 wt.% drastically reduces the athermal martensite share when fast cooling the alloy from room temperature to &#;196 °C followed by up-quenching in water [ 23 ]. However, slow isochronal heating from the cryogenic temperature, or isothermal holding, ideally in the range from &#;80 to &#;40 °C, produces a large volume fraction of isothermal martensite. The activation energy for the formation of this suppressible martensite was found to increase in the range of 0 to 30 kJ/mol depending on the interstitial (C, N) content, however without an apparent influence by the nickel or chrome content [ 23 ]. It can be inferred that the nickel alloy content primarily interferes with the nucleation of athermal martensite.

With increasing nickel alloy content, pearlite and subsequently bainite transformation is progressively suppressed, finally leading to full martensite formation. The maximum transformation rate for combined pearlite and bainite formation decreases with increasing nickel content. In steels with more than 3 wt.% Ni it becomes therefore difficult to separate pearlite from bainite formation. The influence of increasing nickel alloy content up to 9 wt.% Ni on the transformation behavior of low-carbon structural steels is exemplarily demonstrated by the CCT curves in Figure 3

These austenite heterogeneities represented by total grain boundary area and intergranular deformation bands are quantitatively described by the S v parameter, i.e., effective grain boundary area per unit volume. Controlled austenite deformation in combination with accelerated cooling produces very fine-grained ferrite, resulting in high yield strength and low ductile-to-brittle transition temperature upon impact loading. However, particularly in heavy-gauged products, it is often not possible to achieve a high degree of controlled deformation in the central area. Furthermore, the achievable cooling rates are comparably low. Thus, nickel alloying can effectively help to improve mechanical properties in sections where such inferior processing conditions prevail.

The significant extension of the austenite phase field by nickel alloying in combination with its low diffusivity in the austenite lattice delays the austenite-to-ferrite transformation towards lower temperature or extended times. The resulting undercooling causes an increased ferrite nucleation rate and decreased ferrite growth rate leading to a general refinement of the ferrite microstructure. However, Kosazu [ 17 ] has reported that the refining effect of nickel alone and under slow cooling rates is rather modest. Accelerated cooling from temperatures above Ainduces a similar undercooling effect and refinement mechanisms. The nickel alloy effect in combination with that of accelerated cooling can thereby act in synergy ( Figure 2 a) [ 17 18 ]. The excessive application of both parameters, i.e., Ni alloying and accelerated cooling might cause, however, the formation of undesirable coarse bainite structures during the transformation from equiaxed austenite. This is related to the circumstance that only preferential austenite grain boundaries can nucleate ferrite grains. Enhancing the heterogeneity in austenite by controlled deformation below the recrystallization temperature, however, provides a much larger number of ferrite nucleation sites, avoiding the formation of coarse-grained bainitic features ( Figure 2 b).

2.4. Solute Atom Effects of Nickel

) and shear (

) moduli of bcc iron. Ledbetter and Reed [

Ni in wt.% as

E (GPa)

= 205 &#; 1.75&#;

(1)

G (GPa)

= 81 &#; 1.12&#;

(2)

The variation in the lattice parameters of austenite and ferrite, with the addition of an alloying element, strongly depends upon the relative difference in atomic radius between iron and the alloying element: the so-called misfit parameter. Nickel, like manganese and chromium, has an atomic size rather similar to that of iron. Accordingly, it has a small effect on the lattice parameter in bcc iron. Leslie [ 26 ] and, more recently, Hagi [ 27 ] have provided quantitative data for this effect. Host lattice distortion by nickel alloying is accordingly rather small. Nevertheless, solute nickel has a quite significant effect on the Young&#;s () and shear () moduli of bcc iron. Ledbetter and Reed [ 28 ] derived relationships describing the influence of nickel contentin wt.% aswhich can have important consequences, for instance, regarding the fracture behavior of carbon steels. Accordingly, Morris et al. [ 29 ] pointed out that in a Fe-12 wt.%Ni alloy the ideal cleavage and shear strengths are respectively more than 10% and 16% below that of pure Fe, and a ferritic alloy containing more than 32 wt.% Ni would become virtually immune to cleavage.

&#;4 per at.% for nickel which is very similar to that for manganese (1.80&#;2.30 × 10&#;4 1/at.%) and larger than that for chromium (1.20&#;1.90 × 10&#;4 per at.%) [&#;4 per at.% and 3.53&#;4.04 × 10&#;4 per at.% [&#;4 s&#;1) around a 20 MPa yield strength increase per at.% Ni alloy addition can be achieved by solid solution strengthening [31,

The interaction of alloying elements in substitutional solution with defects and interstitials is controlled by the lattice strain caused by the alloying atom. It is reasonable to assume that the solid solution strengthening effect should correlate with the misfit parameter. The normalized lattice distortion is proportional to the concentration of the substitutional element and has been reported to be in the range of 1.65&#;2.31 × 10per at.% for nickel which is very similar to that for manganese (1.80&#;2.30 × 101/at.%) and larger than that for chromium (1.20&#;1.90 × 10per at.%) [ 27 ]. In comparison, larger-sized substitutional atoms such as molybdenum or vanadium cause much bigger misfits, in the range of 5.21&#;8.36 × 10per at.% and 3.53&#;4.04 × 10per at.% [ 27 ]. While chromium, vanadium, and molybdenum fit quite well to a linear correlation between misfit parameter and solid solution strengthening, nickel and manganese strengthen bcc iron more than their misfit parameter suggests [ 30 ]. Under tensile test conditions at ambient temperature (23 °C) and low strain rate (10) around a 20 MPa yield strength increase per at.% Ni alloy addition can be achieved by solid solution strengthening [ 30 ]. On the other hand, binary iron alloys containing Ni (and similarly Mn, Al and Si) exhibit solid solution softening below a temperature of 250 K [ 30 32 ]. Solid solution softening has been attributed to a decrease in the kink-pair nucleation energy with increasing solute content. Following the original suggestion by Weertman, a solute atom misfit center would interact with a dislocation, helping to pull it out of the Peierls energy valley [ 33 34 ]. This interpretation, however, does not assume a direct influence of nickel on the Peierls stress. The experimental results of Hildebrandt and Dickenscheid [ 35 ] prove that the activation energy for slip in Fe&#;Ni alloys is reduced relative to pure iron. This difference increases with the nickel content over the investigated alloy range of up to 3 at.% and also increases the further that temperature decreases below ambient. The authors attributed this behavior to a disturbance in the iron matrix caused by the change of the lattice parameter and elastic modulus due to nickel alloying. Furthermore, it was argued that nickel fills the incomplete 3d electron shell of iron, thus diminishing the influence of covalent bonding so that more active slip systems are available at low temperature. Chen et al. [ 36 ] have presented results on the plastic flow of binary Fe&#;Ni alloys, which are congruent with those reported in [ 35 ]. Their analysis indeed suggests that the Peierls stress decreased from 920 MPa in pure iron to 640 MPa when adding 4 at.%. Ni. The temperature dependence of the strain rate sensitivity exposes a local minimum, which shifts towards lower temperatures with increasing nickel content.

Due to its relatively small misfit parameter, high solubility and slow diffusivity in iron, nickel has only a weak tendency for segregating to grain boundaries. This is contrary to solute atoms inducing a larger misfit such as molybdenum or niobium [ 30 37 ]. The low binding energy of nickel to the boundary results in a comparably small solute drag effect. Empirical work by Lee [ 38 ] indicates that the contribution of nickel in enhancing the activation energy for austenite grain growth is considerably lower than that of molybdenum and also lower than that of chromium. Ab-initio calculations by Hoerner et al. [ 39 ] demonstrate a good correlation between the binding energy of solutes to preferential substitutional sites and the activation energy for austenite growth. Accordingly, nickel has a small impact on obstructing the boundary mobility, which is a major factor regarding austenite recrystallization and grain growth. Research by Jin et al. [ 40 ], using first-principle calculations, suggests that nickel also has a low binding energy to the bcc&#;fcc transformation interface. Hence, its solute drag effect on the phase transformation front is also small. Modelling of alloy effects on ferrite growth proceeding from the transformation of austenite by Zurob et al. [ 41 ] confirmed a weak interaction of nickel with the interface along with the tendency for carbon to repel nickel from the interface.

High temperature plasticity in ferrite relies on the mobility of dislocations and recovery phenomena. Solute alloy atoms can retard the gliding motion of dislocations by dragging. The climbing motion of dislocations is controlled by vacancy flow towards the dislocation core. Yoshida et al. [ 42 ] have argued that the dragging force of solutes on dislocation gliding is determined by the solute misfit parameter and its self-diffusion coefficient. Accordingly, nickel has only a small drag effect which is owed mostly to its limited diffusivity rather than to its small misfit.

The diffusion coefficient of nickel in a highly defective microstructure, however, can be considerably increased as compared to normal lattice diffusion [ 43 ]. Enhanced nickel diffusion has been observed, for instance, after cold-working or neutron irradiation [ 44 ]. Thereby the activation energy for diffusion of nickel in the highly defective steel microstructure is attributed mainly to the activation energy for vacancy migration. In that case, the already weak obstructive effect of solute nickel on dislocation glide or recovery should be further reduced. Enhanced nickel diffusivity is also promoting nickel partitioning during heat treatment in the temperature range of coexisting ferrite and austenite. This is the case when a defect-rich original microstructure such as martensite is tempered in the inter-critical temperature range resulting in a fraction of metastable retained austenite.

2) so that the number of stacking faults is relatively low, explaining its good plasticity. When nickel is alloyed to austenitic steel it also increases the stacking fault energy of the iron matrix. The magnitude of the effect was found to be ~2.8 mJ/m2 per wt.% nickel added [2), chromium (~0.4 mJ/m2) and manganese (~0.75 mJ/m2) per wt.% addition [2, lead to widely dissociated dislocations whose movements are therefore confined to the slip planes, whereas localized cross-slip is extremely difficult. Higher SFE results in a lower barrier to dislocation motion due to the ability of cross slipping without splitting up into partial dislocations. The reduction of Peierls stress, in combination with the increase in dislocation mobility assisted by higher SFE (

Another plasticity-influencing effect of nickel is related to a change in stacking fault energy (SFE). In fcc metals, such as austenite, stacking faults (SF) are interruptions in the perfect (ABCABC) sequential layering of the (111) crystallographic planes. These faults are formed by either dissociation of a perfect dislocation into two partial dislocations, or by emission of partial dislocations from grain boundaries. The stacking fault energy is required for generating this lattice irregularity. Pure nickel metal has a comparably high stacking fault energy (SFE ~128 mJ/m) so that the number of stacking faults is relatively low, explaining its good plasticity. When nickel is alloyed to austenitic steel it also increases the stacking fault energy of the iron matrix. The magnitude of the effect was found to be ~2.8 mJ/mper wt.% nickel added [ 45 ]. This value is larger than that for other alloying elements such as molybdenum (~2.0 mJ/m), chromium (~0.4 mJ/m) and manganese (~0.75 mJ/m) per wt.% addition [ 45 ]. The level of stacking fault energy has an impact on the deformation characteristics of the material such as the work-hardening and softening behavior. Low SFE values, typically less than 45 mJ/m, lead to widely dissociated dislocations whose movements are therefore confined to the slip planes, whereas localized cross-slip is extremely difficult. Higher SFE results in a lower barrier to dislocation motion due to the ability of cross slipping without splitting up into partial dislocations. The reduction of Peierls stress, in combination with the increase in dislocation mobility assisted by higher SFE ( Figure 6 ) in nickel alloyed steels, is the metal-physical origin for improved toughness properties under impact loading (see Section 2.6 ) as well as better ductility in cold-forming processes.

&#;2) and martensite transformation (<18 mJ&#;m&#;2) [

The work-hardening rate is controlled by ease of cross-slip governed by the SFE, which in austenitic stainless steel is influenced by alloying elements [ 46 ]. Regarding carbon steels, the SFE and potential effects by nickel alloying are relevant to hot deformation in the austenite range and the behavior of potentially retained austenite phases in final condition. Strain hardening and softening processes take place alternately during hot rolling of the austenitic microstructure. Softening is caused either by recovery or recrystallization. The increase of the SFE by nickel alloying and the correspondingly improved dislocation mobility should promote recovery as softening mechanism during hot deformation. However, the overall recrystallization kinetics of the steel is not noticeably changed by nickel. Only the solute-drag effect caused by substitutional solutes delays the speed of the necessary grain boundary movement for a given recrystallization driving force. Retained austenite in carbon steel deforms by various secondary plasticity mechanisms, such as mechanical twinning (18 &#; γ &#; 45 mJ&#;m) and martensite transformation (<18 mJ&#;m) [ 47 ]. In that respect, the amount of nickel partitioned to the retained austenite might influence the mechanical properties of this phase.

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